ASM Handbook: Volume 6: Welding, Brazing, and Soldering (Asm Handbook)

VOLUME ASM INTERNATIONAL ® PUBLICATION INFORMATION AND CONTRIBUTORS WELDING, BRAZING, AND SOLDERING WAS PUBLISHED I...

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VOLUME

ASM INTERNATIONAL

®

PUBLICATION INFORMATION AND CONTRIBUTORS

WELDING, BRAZING, AND SOLDERING WAS PUBLISHED IN 1993 AS VOLUME 6 OF THE ASM HANDBOOK. THE VOLUME WAS PREPARED UNDER THE DIRECTION OF THE ASM HANDBOOK COMMITTEE. VOLUME CHAIRMEN THE VOLUME CHAIRMEN WERE DAVID LEROY OLSON, THOMAS A. SIEWERT, STEPHEN LIU, AND GLEN R. EDWARDS. AUTHORS • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • • •

LAMET UFRGS BRUNO L. ALIA RICHARD L. ALLEY AMERICAN WELDING SOCIETY WILLIAM R. APBLETT, JR. WILLIAM A. BAESLACK III THE OHIO STATE UNIVERSITY WILLIAM BALLIS COLUMBIA GAS OF OHIO CLIFF C. BAMPTON ROCKWELL INTERNATIONAL SCIENCE CENTER PROBAL BANERJEE AUBURN UNIVERSITY JOHN G. BANKER EXPLOSIVE FABRICATORS INC. ROBERT G. BARTIFAY ALUMINUM COMPANY OF AMERICA ROY I. BATISTA ROY E. BEAL AMALGAMATED TECHNOLOGIES INC. RAYMOND E. BOHLMANN MCDONNELL AIRCRAFT COMPANY SÉRGIO D. BRANDI ESCOLA POLITECNICA DA USP JOHN A. BROOKS SANDIA NATIONAL LABORATORIES DONALD W. BUCHOLZ IBM FEDERAL SYSTEMS CORPORATION PAUL BURGARDT EG&G ROCKY FLATS PLANT ROGER A. BUSHEY THE ESAB GROUP INC. CHRIS CABLE FEIN POWER TOOL RICHARD D. CAMPBELL JOINING SERVICES INC. HOWARD CARY HOBART BROTHERS COMPANY HARVEY CASTNER EDISON WELDING INSTITUTE ALLEN CEDILOTE INDUSTRIAL TESTING LABORATORY SERVICES HARRY A. CHAMBERS TRW NELSON STUD WELDING C. CHRIS CHEN MICROALLOYING INTERNATIONAL INC. SHAOFENG CHEN AUBURN UNIVERSITY SHAO-PING CHEN LOS ALAMOS NATIONAL LABORATORY BRYAN A. CHIN AUBURN UNIVERSITY MICHAEL J. CIESLAK SANDIA NATIONAL LABORATORIES RODGER E. COOK THE WILKINSON COMPANY STEPHEN A. COUGHLIN ACF INDUSTRIES INC. MARK COWELL METCAL INC. RICHARD S. CREMISIO RESCORP INTERNATIONAL INC. CARL E. CROSS CRAIG DALLAM THE LINCOLN ELECTRIC COMPANY BRIAN DAMKROGER SANDIA NATIONAL LABORATORIES JOSEPH R. DAVIS DAVIS AND ASSOCIATES

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JANET DEVINE SONOBOND ULTRASONICS PAUL B. DICKERSON RAY DIXON LOS ALAMOS NATIONAL LABORATORY SUE DUNKERTON THE WELDING INSTITUTE KEVIN DUNN TEXAS INSTRUMENTS INC. CHUCK DVORAK UNI-HYDRO, INC. JIM DVORAK UNI-HYDRO, INC. ROBERT J. DYBAS GENERAL ELECTRIC COMPANY THOMAS W. EAGAR MASSACHUSETTS INSTITUTE OF TECHNOLOGY GLEN R. EDWARDS COLORADO SCHOOL OF MINES GRAHAM R. EDWARDS THE WELDING INSTITUTE W.H. ELLIOTT, JR. OAK RIDGE NATIONAL LABORATORY JOHN W. ELMER LAWRENCE LIVERMORE NATIONAL LABORATORY STEVEN C. ERNST EASTMAN CHEMICAL COMPANY WILLIAM FARRELL FERRANTI-SCIAKY COMPANY JOEL G. FELDSTEIN FOSTER WHEELER ENERGY CORPORATION DAVID A. FLEMING COLORADO SCHOOL OF MINES JAMES A. FORSTER TEXAS INSTRUMENTS INC. MICHAEL D. FREDERICKSON ELECTRONICS MANUFACTURING PRODUCTIVITY FACILITY EDWARD FRIEDMAN WESTINGHOUSE ELECTRIC CORPORATION R.H. FROST COLORADO SCHOOL OF MINES CHARLES E. FUERSTENAU LUCAS-MILHAUPT INC. EDWARD B. GEMPLER STANLEY S. GLICKSTEIN WESTINGHOUSE ELECTRIC CORPORATION JOHN A. GOLDAK CARLETON UNIVERSITY ROBIN GORDON EDISON WELDING INSTITUTE JERRY E. GOULD EDISON WELDING INSTITUTE JOHN B. GREAVES, JR. ELECTRONICS MANUFACTURING PRODUCTIVITY FACILITY F. JAMES GRIST JOHN F. GRUBB ALLEGHENY LUDLUM STEEL MAOSHI GU CARLETON UNIVERSITY IAN D. HARRIS EDISON WELDING INSTITUTE L.J. HART-SMITH DOUGLAS AIRCRAFT COMPANY DAN HAUSER EDISON WELDING INSTITUTE C.R. HEIPLE METALLURGICAL CONSULTANT HERBERT HERMAN STATE UNIVERSITY OF NEW YORK G. KEN HICKEN SANDIA NATIONAL LABORATORIES EVAN B. HINSHAW INCO ALLOYS INTERNATIONAL INC. D. BRUCE HOLLIDAY WESTINGHOUSE MARINE DIVISION S. IBARRA AMOCO CORPORATION J. ERNESTO INDACOCHEA UNIVERSITY OF ILLINOIS AT CHICAGO SUNIL JHA TEXAS INSTRUMENTS INC. JERALD E. JONES COLORADO SCHOOL OF MINES RAYMOND H. JUERS NAVAL SURFACE WARFARE CENTER WILLIAM R. KANNE, JR. WESTINGHOUSE SAVANNAH RIVER COMPANY MICHAEL J. KARAGOULIS GENERAL MOTORS CORPORATION MICHAEL KARAVOLIS TEXAS INSTRUMENTS INC. LENNART KARLSSON LULEÅ UNIVERSITY OF TECHNOLOGY MICHAEL E. KASSNER OREGON STATE UNIVERSITY DOUG D. KAUTZ LAWRENCE LIVERMORE NATIONAL LABORATORY W. DANIEL KAY WALL COLMONOY CORPORATION JAMES F. KEY IDAHO NATIONAL ENGINEERING LABORATORY H.-E. KIM SEOUL NATIONAL UNIVERSITY

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SAMUEL D. KISER INCO ALLOYS INTERNATIONAL INC. MARVIN L. KOHN FMC CORPORATION DAMIAN J. KOTECKI THE LINCOLN ELECTRIC COMPANY KENNETH KRYSIAC HERCULES INC. CHUCK LANDRY THERMAL DYNAMICS CHARLES LANE DURALCAN H.J. LATIMER TAYLOR-WINFIELD CORPORATION GLEN S. LAWRENCE FERRANTI-SCIAKY COMPANY KARL LAZAR WERNER LEHRHEUER FORSCHUNGSZENTRUM JÜLICH GMBH ALEXANDER LESNEWICH J.F. LIBSCH LEPEL CORPORATION TOM LIENERT THE OHIO STATE UNIVERSITY ALLEN C. LINGENFELTER LAWRENCE LIVERMORE NATIONAL LABORATORY DALE L. LINMAN CENTECH CORPORATION VONNE LINSE EDISON WELDING INSTITUTE JOHN C. LIPPOLD EDISON WELDING INSTITUTE JIAYAN LIU AUBURN UNIVERSITY STEPHEN LIU COLORADO SCHOOL OF MINES MATTHEW J. LUCAS, JR. GENERAL ELECTRIC COMPANY KEVIN A. LYTTLE PRAXAIR INC. KIM MAHIN SANDIA NATIONAL LABORATORIES MURRAY W. MAHONEY ROCKWELL INTERNATIONAL SCIENCE CENTER DARRELL MANENTE VAC-AERO INTERNATIONAL INC. RICHARD P. MARTUKANITZ PENNSYLVANIA STATE UNIVERSITY KOICHI MASUBUCHI MASSACHUSETTS INSTITUTE OF TECHNOLOGY DAVID K. MATLOCK COLORADO SCHOOL OF MINES R.B. MATTESON TAYLOR-WINFIELD CORPORATION STEVEN J. MATTHEWS HAYNES INTERNATIONAL INC. JYOTI MAZUMDER UNIVERSITY OF ILLINOIS AT URBANA-CHAMPAIGN C.N. MCCOWAN NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY KRIS MEEKINS LONG MANUFACTURING LTD. GREGORY MELEKIAN GENERAL MOTORS CORPORATION ANTHONY R. MELLINI, SR. MELLINI AND ASSOCIATES INC. DAVID W. MEYER THE ESAB GROUP INC. JULE MILLER HOWARD MIZUHARA WESGO INC. ARTHUR G. MOORHEAD OAK RIDGE NATIONAL LABORATORY MILO NANCE MARTIN MARIETTA ASTRONAUTICS GROUP E.D. NICHOLAS THE WELDING INSTITUTE DAVID NOBLE ARCO EXPLORATION AND PRODUCTION TECHNOLOGY THOMAS NORTH UNIVERSITY OF TORONTO DAVID B. O'DONNELL INCO ALLOYS INTERNATIONAL INC. JONATHAN S. OGBORN THE LINCOLN ELECTRIC COMPANY DAVID L. OLSON COLORADO SCHOOL OF MINES TOSHI OYAMA WESGO INC. R. ALAN PATTERSON LOS ALAMOS NATIONAL LABORATORY LARRY PERKINS WRIGHT LABORATORY DARYL PETER DARYL PETER AND ASSOCIATES MANFRED PETRI GERHARD PETRI GMBH & CO. KG DAVID H. PHILLIPS EDISON WELDING INSTITUTE ABE POLLACK MICROALLOYING INTERNATIONAL INC. BARRY POLLARD ANATOL RABINKIN ALLIEDSIGNAL AMORPHOUS METALS

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GEETHA RAMARATHNAM UNIVERSITY OF TORONTO EDWARD G. REINEKE EXPLOSIVE FABRICATORS INC. JULIAN ROBERTS THERMATOOL CORPORATION M. NED ROGERS BATESVILLE CASKET COMPANY J.R. ROPER EG&G ROCKY FLATS PLANT ROBERT S. ROSEN LAWRENCE LIVERMORE NATIONAL LABORATORY JAMES E. ROTH JAMES E. ROTH INC. WILLIAM J. RUPRECHT GENERAL ELECTRIC COMPANY K. SAMPATH CONCURRENT TECHNOLOGIES CORPORATION BERNARD E. SCHALTENBRAND ALUMINUM COMPANY OF AMERICA BERNARD SCHWARTZ NORFOLK SOUTHERN CORPORATION MEL M. SCHWARTZ SIKORSKY AIRCRAFT ANN SEVERIN LUCAS-MILHAUPT INC. THOMAS A. SIEWERT NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY HERSCHEL SMARTT IDAHO NATIONAL ENGINEERING LABORATORY RONALD B. SMITH ALLOY RODS CORPORATION WARREN F. SMITH THERMATOOL CORPORATION LANCE R. SOISSON WELDING CONSULTANTS INC. HARVEY D. SOLOMON GENERAL ELECTRIC COMPANY BRUCE R. SOMERS LEHIGH UNIVERSITY ROBERT E. SOMERS SOMERS CONSULTANTS ROGER K. STEELE AAR TECHNICAL CENTER FRANK STEIN TAYLOR-WINFIELD CORPORATION TIM STOTLER EDISON WELDING INSTITUTE ROBERT L. STROHL TWECO/ARCAIR ROBERT A. SULIT SULIT ENGINEERING VERN SUTTER AMERICAN WELDING INSTITUTE W.T. TACK MARTIN MARIETTA R. DAVID THOMAS, JR. R.D. THOMAS & COMPANY KARL THOMAS TECHNISCHE UNIVERSITÄT, BRAUNSCHWEIG RAYMOND G. THOMPSON UNIVERSITY OF ALABAMA AT BIRMINGHAM DONALD J. TILLACK D.J. TILLACK & ASSOCIATES CHON L. TSAI THE OHIO STATE UNIVERSITY SCHILLINGS TSANG EG&G ROCKY FLATS PLANT HENDRIKUS H. VANDERVELDT AMERICAN WELDING INSTITUTE RICCARDO VANZETTI VANZETTI SYSTEMS INC. PAUL T. VIANCO SANDIA NATIONAL LABORATORIES P. RAVI VISHNU LULEÅ UNIVERSITY OF TECHNOLOGY MARY B. VOLLARO UNIVERSITY OF CONNECTICUT A. WAHID COLORADO SCHOOL OF MINES DANIEL W. WALSH CALIFORNIA POLYTECHNIC STATE UNIVERSITY R. TERRENCE WEBSTER CONSULTANT JOHN R. WHALEN CONTOUR SAWS INC. NEVILLE T. WILLIAMS BRITISH STEEL FRED J. WINSOR WELDING CONSULTANT R. XU UNIVERSITY OF ILLINOIS AT CHICAGO XIAOSHU XU AMERICAN WELDING INSTITUTE PHILIP M. ZARROW SYNERGISTEK ASSOCIATES

REVIEWERS • • •

YONI ADONYI U.S. STEEL TECHNICAL CENTER RICHARD L. ALLEY AMERICAN WELDING SOCIETY BERNARD ALTSHULLER ALCAN INTERNATIONAL LTD.

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TED L. ANDERSON TEXAS A&M UNIVERSITY LLOYD ANDERSON MARION-INDRESCO INC. FRANK G. ARMAO ALCOA TECHNICAL CENTER DANIEL ARTHUR TELEDYNE MCKAY RICHARD E. AVERY NICKEL DEVELOPMENT INSTITUTE R.F. BACON TECUMSEH PRODUCTS COMPANY WALLY G. BADER WILLIAM A. BAESLACK III THE OHIO STATE UNIVERSITY CLIFF C. BAMPTON ROCKWELL INTERNATIONAL SCIENCE CENTER JOHN G. BANKER EXPLOSIVE FABRICATORS INC. GEORGE C. BARNES ROBERT G. BARTIFAY ALUMINUM COMPANY OF AMERICA ROY E. BEAL AMALGAMATED TECHNOLOGIES INC. GARY BECKA ALLIEDSIGNAL AEROSPACE COMPANY DAN BEESON EXXON PRODUCTION MALAYSIA DAVID M. BENETEAU CENTERLINE (WINDSOR) LTD. CHRISTOPHER C. BERNDT THE THERMAL SPRAY LABORATORY SURENDRA BHARGAVA GENERAL MOTORS INC. NORMAN C. BINKLEY EDISON WELDING INSTITUTE ROBERT A. BISHEL INCO ALLOYS INTERNATIONAL INC. R.A. BLACK BLACKS EQUIPMENT LTD. OMER W. BLODGETT THE LINCOLN ELECTRIC COMPANY RICHARD A. BRAINARD GENERAL DYNAMICS LAND SYSTEMS DIVISION GLENN H. BRAVE ASSOCIATION OF AMERICAN RAILROADS ROBERT S. BROWN CARPENTER TECHNOLOGY CORPORATION WILLIAM A. BRUCE EDISON WELDING INSTITUTE CHUCK CADDEN GENERAL MOTORS HARVEY R. CASTNER EDISON WELDING INSTITUTE ALLEN B. CEDILOTE INDUSTRIAL TESTING LABORATORY SERVICES CORPORATION KENNETH D. CHALLENGER SAN JOSE STATE UNIVERSITY P.R. CHIDAMBARAM COLORADO SCHOOL OF MINES BOB CHRISTOFFEL ROBIN CHURCHILL ESCO CORPORATION MICHAEL J. CIESLAK SANDIA NATIONAL LABORATORIES BRADLEY A. CLEVELAND MTS SYSTEMS CORPORATION NANCY C. COLE OAK RIDGE NATIONAL LABORATORY HAROLD R. CONAWAY ROCKWELL INTERNATIONAL RICHARD B. CORBIT GENERAL PUBLIC UTILITIES NUCLEAR CORPORATION MARK COWELL METCAL INC. NORM COX RESEARCH INC. JOHN A. CRAWFORD NAVAL SURFACE WARFARE CENTER DENNIS D. CROCKETT THE LINCOLN ELECTRIC COMPANY CARL E. CROSS NARENDRA B. DAHOTRE UNIVERSITY OF TENNESSEE SPACE INSTITUTE T. DEBROY PENNSYLVANIA STATE UNIVERSITY JOSEPH DEVITO THE ESAB GROUP INC. JOHN A. DEVORE GENERAL ELECTRIC COMPANY PAUL B. DICKERSON RAY DIXON LOS ALAMOS NATIONAL LABORATORY KARL E. DORSCHU WELDRING COMPANY INC. ROBERT J. DYBAS GENERAL ELECTRIC COMPANY THOMAS W. EAGAR MASSACHUSETTS INSTITUTE OF TECHNOLOGY BRUCE J. EBERHARD WESTINGHOUSE SAVANNAH RIVER COMPANY GLEN R. EDWARDS COLORADO SCHOOL OF MINES

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JOHN W. ELMER LAWRENCE LIVERMORE NATIONAL LABORATORY WERNER ENGELMAIER ENGELMAIER ASSOCIATES INC. CHRIS ENGLISH GE AIRCRAFT ENGINES ROBERT G. FAIRBANKS SCARROTT METALLURGICAL COMPANY HOWARD N. FARMER CONSULTANT DAVID A. FLEMING COLORADO SCHOOL OF MINES ROBERT FOLEY COLORADO SCHOOL OF MINES BOBBY FOLKENING FMC GROUND SYSTEMS DIVISION DARREL FREAR SANDIA NATIONAL LABORATORIES MICHAEL D. FREDERICKSON ELECTRONICS MANUFACTURING PRODUCTIVITY FACILITY EUGENE R. FREULER SOUDRONIC NEFTENBACH AG STEVEN A. GEDEON WELDING INSTITUTE OF CANADA BOB GIBBONS PLS MATERIALS INC. PAUL S. GILMAN ALLIEDSIGNAL INC. STANLEY S. GLICKSTEIN WESTINGHOUSE ELECTRIC CORPORATION JOHN A. GOLDAK CARLETON UNIVERSITY CARL GRAF EDISON WELDING INSTITUTE WILLIAM L. GREEN THE OHIO STATE UNIVERSITY CHUCK GREGOIRE NATIONAL STEEL CORPORATION ROBERT A. GRIMM EDISON WELDING INSTITUTE BRIAN GRINSELL THOMPSON WELDING INC. ROBIN GROSS-GOURLEY WESTINGHOUSE JOHN F. GRUBB ALLEGHENY LUDLUM STEEL BOB GUNOW, JR. VAC-MET INC. C. HOWARD HAMILTON WASHINGTON STATE UNIVERSITY JAMES R. HANNAHS PMI FOOD EQUIPMENT GROUP FRANK HANNEY ABCO WELDING & INDUSTRIAL SUPPLY INC. DAVID E. HARDT MASSACHUSETTS INSTITUTE OF TECHNOLOGY IAN D. HARRIS EDISON WELDING INSTITUTE MARK J. HATZENBELLER KRUEGER INTERNATIONAL DAN HAUSER EDISON WELDING INSTITUTE C.R. HEIPLE METALLURGICAL CONSULTANT J.S. HETHERINGTON HETHERINGTON INC. BARRY S. HEUER NOOTER CORPORATION ROGER B. HIRSCH UNITROL ELECTRONICS INC. TIM P. HIRTHE LUCAS-MILHAUPT HUGH B. HIX INTERNATIONAL EXPLOSIVE METALWORKING ASSOCIATION F. GALEN HODGE HAYNES INTERNATIONAL INC. RICHARD L. HOLDREN WELDING CONSULTANTS INC. ALAN B. HOPPER ROBERTSHAW TENNESSEE DIVISION CHARLES HUTCHINS C. HUTCHINS AND ASSOCIATES JENNIE S. HWANG IEM-FUSION INC. S. IBARRA AMOCO CORPORATION J. ERNESTO INDACOCHEA UNIVERSITY OF ILLINOIS AT CHICAGO GARY IRONS HOBART TAFA TECHNOLOGIES INC. JAMES R. JACHNA MODINE MANUFACTURING COMPANY ROBERT G. JAITE WOLFENDEN INDUSTRIES INC. JOHN C. JENKINS CONSULTANT KATHI JOHNSON HEXACON ELECTRIC COMPANY WILLIAM R. JONES VACUUM FURNACE SYSTEMS CORPORATION ROBERT W. JUD CHRYSLER CORPORATION WILLIAM F. KAUKLER UNIVERSITY OF ALABAMA IN HUNTSVILLE DOUG D. KAUTZ LAWRENCE LIVERMORE NATIONAL LABORATORY

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W. DANIEL KAY WALL COLMONOY CORPORATION JACQUE KENNEDY WESTINGHOUSE JAMES F. KING OAK RIDGE NATIONAL LABORATORY ANDREW G. KIRETA COPPER DEVELOPMENT ASSOCIATION INC. SAMUEL D. KISER INCO ALLOYS INTERNATIONAL INC. JOSEPH H. KISSEL ITT STANDARD FRED KOHLER CONSULTANT M.L. KOHN FMC CORPORATION DAMIAN J. KOTECKI THE LINCOLN ELECTRIC COMPANY SINDO KOU UNIVERSITY OF WISCONSIN-MADISON CURTIS W. KOVACH TECHNICAL MARKETING RESOURCES INC. LAWRENCE S. KRAMER MARTIN MARIETTA LABORATORIES RAYMOND B. KRIEGER AMERICAN CYANAMID COMPANY KENNETH KRYSIAC HERCULES INC. DANIEL KURUZAR MANUFACTURING TECHNOLOGY INC. RICHARD A. LAFAVE ELLIOTT COMPANY FRANK B. LAKE THE ESAB GROUP INC. JOHN D. LANDES UNIVERSITY OF TENNESSEE WERNER LEHRHEUER FORSCHUNGSZENTRUM JÜLICH GMBH J.F. LIBSCH LEPEL CORPORATION VONNE LINSE EDISON WELDING INSTITUTE JOHN C. LIPPOLD EDISON WELDING INSTITUTE STEPHEN LIU COLORADO SCHOOL OF MINES RONALD LOEHMAN ADVANCED MATERIALS LABORATORY PAUL T. LOVEJOY ALLEGHENY LUDLUM STEEL GEORGE LUCEY U.S. ARMY LABORATORY COMMAND KEVIN A. LYTTLE PRAXAIR INC. COLIN MACKAY MICROELECTRONICS AND COMPUTER TECHNOLOGY CORPORATION MICHAEL C. MAGUIRE SANDIA NATIONAL LABORATORIES KIM W. MAHIN SANDIA NATIONAL LABORATORIES WILLIAM E. MANCINI DUPONT DARRELL MANENTE VAC-AERO INTERNATIONAL INC. AUGUST F. MANZ A.F. MANZ ASSOCIATES RICHARD P. MARTUKANITZ PENNSYLVANIA STATE UNIVERSITY KOICHI MASUBUCHI MASSACHUSETTS INSTITUTE OF TECHNOLOGY STEVEN J. MATTHEWS HAYNES INTERNATIONAL JYOTI MAZUMDER UNIVERSITY OF ILLINOIS AT URBANA-CHAMPAIGN JIM MCMAHON DOALL COMPANY ALAN MEIER COLORADO SCHOOL OF MINES STANLEY MERRICK TELEDYNE MCKAY ROBERT W. MESSLER, JR. RENSSELAER POLYTECHNIC INSTITUTE E.A. METZBOWER U.S. NAVAL RESEARCH LABORATORY JOEL MILANO DAVID TAYLOR MODEL BASIN ROBERT A. MILLER SULZER PLASMA TECHNIK INC. HERBERT W. MISHLER EDISON WELDING INSTITUTE BRAJENDRA MISHRA COLORADO SCHOOL OF MINES HOWARD MIZUHARA WESGO INC. RICHARD MONTANA MID-FLORIDA TECHNICAL INSTITUTE JERRY MOODY WORLD WIDE WELDING RICHARD A. MORRIS NAVAL SURFACE WARFARE CENTER P.J. MUDGE THE WELDING INSTITUTE AMIYA MUKHERJEE UNIVERSITY OF CALIFORNIA BILL MYERS DRESSER-RAND INC.

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ERNEST F. NIPPES CONSULTANT DONG WON OH COLORADO SCHOOL OF MINES DAVID L. OLSON COLORADO SCHOOL OF MINES EDGAR D. OPPENHEIMER CONSULTING ENGINEER CARMEN PAPONETTI HI TECMETAL GROUP INC. MADHU PAREKH HOBART BROTHERS COMPANY SUBHASH R. PATI INTERNATIONAL PAPER COMPANY R. ALAN PATTERSON LOS ALAMOS NATIONAL LABORATORIES CHARLES C. PEASE CP METALLURGICAL ROBERT LEON PEASLEE WALL COLMONOY CORPORATION DARYL PETER DARYL PETER & ASSOCIATES LORENZ PFEIFER JOHN F. PFLZNIENSKI KOLENE CORPORATION DAVID H. PHILLIPS EDISON WELDING INSTITUTE EARL W. PICKERING, JR. CONSULTANT E.R. PIERRE CONSULTING WELDING ADVISOR JOHN PILLING MICHIGAN TECHNOLOGICAL UNIVERSITY ABE POLLACK MICROALLOYING INTERNATIONAL INC. BARRY POLLARD ALEXANDRE M. POPE COLORADO SCHOOL OF MINES JEFFREY W. POST J.W. POST & ASSOCIATES INC. TERRY PROFUGHI HI TECMETAL GROUP INC. ANATOL RABINKIN ALLIEDSIGNAL AMORPHOUS METALS JIM D. RABY SOLDERING TECH INTERNATIONAL TED RENSHAW CONSULTANT THERESA ROBERTS SIKAMA INTERNATIONAL DAVID E. ROBERTSON PACE INC. CHARLES ROBINO SANDIA NATIONAL LABORATORIES M.N. ROGERS ABB POWER T&D COMPANY INC. J.R. ROPER EG&G ROCKY FLATS PLANT N.V. ROSS AJAX MAGNETHERMIC DIETRICH K. ROTH ROMAN MANUFACTURING INC. JOHN RUFFING 3M FLUIDS LABORATORY WAYNE D. RUPERT ENGLEHARD CORPORATION J.D. RUSSELL THE WELDING INSTITUTE C.O. RUUD PENNSYLVANIA STATE UNIVERSITY EDMUND F. RYBICKI UNIVERSITY OF TULSA JONATHAN T. SALKIN ARC APPLICATIONS INC. MEL M. SCHWARTZ SIKORSKY AIRCRAFT JOE L. SCOTT DEVASCO INTERNATIONAL INC. ALAN P. SEIDLER RMI TITANIUM COMPANY OSCAR W. SETH CHICAGO BRIDGE & IRON COMPANY ANN SEVERIN LUCAS-MILHAUPT INC. LEWIS E. SHOEMAKER INCO ALLOYS INTERNATIONAL INC. LYNN E. SHOWALTER NEWPORT NEWS SHIPBUILDING THOMAS A. SIEWERT NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY ALLEN W. SINDEL BEGEMANN HEAVY INDUSTRIES INC. MICHAEL H. SKILLINGBERG REYNOLDS METALS COMPANY GERALD M. SLAUGHTER OAK RIDGE NATIONAL LABORATORY HERSCHEL SMARTT IDAHO NATIONAL ENGINEERING LABORATORY JAMES P. SNYDER II BETHLEHEM STEEL CORPORATION LANCE R. SOISSON WELDING CONSULTANTS INC. HARVEY D. SOLOMON GENERAL ELECTRIC BRUCE R. SOMERS LEHIGH UNIVERSITY

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NARASI SRIDHAR SOUTHWEST RESEARCH INSTITUTE BOB STANLEY NATIONAL TRAINING FUND ROGER K. STEELE ASSOCIATION OF AMERICAN RAILROADS ARCHIE STEVENSON MAGNESIUM ELEKRON INC. VIJAY K. STOKES GENERAL ELECTRIC TIM STOTLER EDISON WELDING INSTITUTE M.A. STREICHER CONSULTANT ROBERT L. STROHL TWECO/ARCAIR LAWRENCE STRYKER ALTECH INTERNATIONAL MARK TARBY WALL COLMONOY CORPORATION CLAY TAYLOR MERRICK AND COMPANY J.R. TERRILL CONSULTANT RAYMOND G. THOMPSON UNIVERSITY OF ALABAMA AT BIRMINGHAM J.S. THROWER GENERAL ELECTRIC POWER GENERATION DONALD J. TILLACK D.J. TILLACK & ASSOCIATES FELIX TOMEI TRUMPF INC. CHON L. TSAI THE OHIO STATE UNIVERSITY SCHILLINGS TSANG EG&G ROCKY FLATS PLANT M. NASIM UDDIN THYSSEN STEEL GROUP ELMAR UPITIS CBI TECHNICAL SERVICES COMPANY JAMES VAN DEN AVYLE SANDI NATIONAL LABORATORIES CLARENCE VAN DYKE LUCAS-MIHAUPT INC. HENDRIKUS H. VANDERVELDT AMERICAN WELDING INSTITUTE DAVID B. VEVERKA EDISON WELDING INSTITUTE PAUL T. VIANCO SANDIA NATIONAL LABORATORIES ROBERT G. VOLLMER R. WALLACH UNIVERSITY OF CAMBRIDGE SANDRA J. WALMSLEY WESTINGHOUSE ELECTRIC CORPORATION RICHARD A. WATSON THE P&LE CAR COMPANY CHRIS WEHLUS GENERAL MOTORS C.E.T. WHITE INDIUM CORPORATION OF AMERICA ROGER N. WILD ELLIOTT WILLNER LOCKHEED MISSILES & SPACE COMPANY RICHARD WILSON HOUSTON LIGHTING AND POWER COMPANY W.L. WINTERBOTTOM FORD MOTOR COMPANY A.P. WOODFIELD GENERAL ELECTRIC AIRCRAFT ENGINES JAMES B.C. WU STOODY COMPANY THOMAS ZACHARIA OAK RIDGE NATIONAL LABORATORY

FOREWORD COVERAGE OF JOINING TECHNOLOGIES IN THE ASM HANDBOOK HAS GROWN DRAMATICALLY OVER THE YEARS. A SHORT CHAPTER ON WELDING--EQUAL IN SIZE TO ABOUT 5 PAGES OF TODAY'S ASM HANDBOOK--APPEARED IN THE 1933 EDITION OF THE NATIONAL METALS HANDBOOK PUBLISHED BY THE AMERICAN SOCIETY OF STEEL TREATERS, ASM'S PREDECESSOR. THAT MATERIAL WAS EXPANDED TO 13 PAGES IN THE CLASSIC 1948 EDITION OF METALS HANDBOOK. THE FIRST FULL VOLUME ON WELDING AND BRAZING IN THE SERIES APPEARED IN 1971, WITH PUBLICATION OF VOLUME 6 OF THE 8TH EDITION OF METALS HANDBOOK. VOLUME 6 OF THE 9TH EDITION, PUBLISHED IN 1983, WAS EXPANDED TO INCLUDE COVERAGE OF SOLDERING. THE NEW VOLUME 6 OF THE ASM HANDBOOK BUILDS ON THE PROUD TRADITION ESTABLISHED BY THESE PREVIOUS VOLUMES, BUT IT ALSO REPRESENTS A BOLD NEW STEP FOR THE SERIES. THE HANDBOOK HAS NOT ONLY BEEN REVISED, BUT ALSO ENTIRELY

REFORMATTED TO MEET THE NEEDS OF TODAY'S MATERIALS COMMUNITY. OVER 90% OF THE ARTICLES IN THIS VOLUME ARE BRAND-NEW, AND THE REMAINDER HAVE BEEN SUBSTANTIALLY REVISED. MORE SPACE HAS BEEN DEVOTED TO COVERAGE OF SOLIDSTATE WELDING PROCESSES, MATERIALS SELECTION FOR JOINED ASSEMBLIES, WELDING IN SPECIAL ENVIRONMENTS, QUALITY CONTROL, AND MODELING OF JOINING PROCESSES, TO NAME BUT A FEW. INFORMATION ALSO HAS BEEN ADDED FOR THE FIRST TIME ABOUT JOINING OF SELECTED NONMETALLIC MATERIALS. WHILE A DELIBERATE ATTEMPT HAS BEEN MADE TO INCREASE THE AMOUNT OF CUTTINGEDGE INFORMATION PROVIDED, THE ORGANIZERS HAVE WORKED HARD TO ENSURE THAT THE HEART OF THE BOOK REMAINS PRACTICAL INFORMATION ABOUT JOINING PROCESSES, APPLICATIONS, AND MATERIALS WELDABILITY--THE TYPE OF INFORMATION THAT IS THE HALLMARK OF THE ASM HANDBOOK SERIES. PUTTING TOGETHER A VOLUME OF THIS MAGNITUDE IS AN ENORMOUS EFFORT AND COULD NOT HAVE BEEN ACCOMPLISHED WITHOUT THE DEDICATED AND TIRELESS EFFORTS OF THE VOLUME CHAIRPERSONS: DAVID L. OLSON, THOMAS A. SIEWERT, STEPHEN LIU, AND GLEN R. EDWARDS. SPECIAL THANKS ARE ALSO DUE TO THE SECTION CHAIRPERSONS, TO THE MEMBERS OF THE ASM HANDBOOK COMMITTEE, AND TO THE ASM EDITORIAL STAFF. WE ARE ESPECIALLY GRATEFUL TO THE OVER 400 AUTHORS AND REVIEWERS WHO HAVE CONTRIBUTED THEIR TIME AND EXPERTISE IN ORDER TO MAKE THIS HANDBOOK A TRULY OUTSTANDING INFORMATION RESOURCE. EDWARD H. KOTTCAMP, JR. PRESIDENT ASM INTERNATIONAL EDWARD L. LANGER MANAGING DIRECTOR ASM INTERNATIONAL PREFACE THE ASM HANDBOOK, VOLUME 6, WELDING, BRAZING, AND SOLDERING, HAS BEEN ORGANIZED INTO A UNIQUE FORMAT THAT WE BELIEVE WILL PROVIDE HANDBOOK USERS WITH READY ACCESS TO NEEDED MATERIALS-ORIENTED JOINING INFORMATION AT A MINIMAL LEVEL OF FRUSTRATION AND STUDY TIME. WHEN WE DEVELOPED THE ORGANIZATIONAL STRUCTURE FOR THIS VOLUME, WE RECOGNIZED THAT ENGINEERS, TECHNICIANS, RESEARCHERS, DESIGNERS, STUDENTS, AND TEACHERS DO NOT SEEK OUT JOINING INFORMATION WITH THE SAME LEVEL OF UNDERSTANDING, OR WITH THE SAME NEEDS. THEREFORE, WE ESTABLISHED DISTINCT SECTIONS THAT WERE INTENDED TO MEET THE SPECIFIC NEEDS OF PARTICULAR USERS. THE EXPERIENCED JOINING SPECIALIST CAN TURN TO THE SECTION "CONSUMABLE SELECTION, PROCEDURE DEVELOPMENT, AND PRACTICE CONSIDERATIONS" AND FIND DETAILED JOINING MATERIALS DATA ON A WELL-DEFINED PROBLEM. THIS HANDBOOK ALSO PROVIDES GUIDANCE FOR THOSE WHO NOT ONLY MUST SPECIFY THE JOINING PRACTICE, BUT ALSO THE MATERIALS TO BE JOINED. THE SECTION "MATERIALS SELECTION FOR JOINED ASSEMBLIES" CONTAINS COMPREHENSIVE INFORMATION ABOUT THE PROPERTIES, APPLICATIONS, AND WELDABILITIES OF THE MAJOR CLASSES OF STRUCTURAL MATERIALS. TOGETHER, THESE TWO MAJOR SECTIONS OF THE HANDBOOK SHOULD PROVIDE AN ENGINEER ASSIGNED A LOOSELY DEFINED DESIGN PROBLEM WITH THE MEANS TO MAKE INTELLIGENT CHOICES FOR COMPLETING AN ASSEMBLY. FREQUENTLY, TECHNOLOGISTS ARE CALLED UPON TO INITIATE AND ADOPT WELDING PROCESSES WITHOUT IN-DEPTH KNOWLEDGE OF THESE PROCESSES OR THE SCIENTIFIC

PRINCIPLES THAT IMPACT THE PROPERTIES AND PERFORMANCE OF WELDMENTS. THE SECTIONS "FUNDAMENTALS OF JOINING" AND "JOINING PROCESSES" ARE DESIGNED TO MEET THE NEEDS OF THESE USERS, OR ANYONE WHO NEEDS BASIC BACKGROUND INFORMATION ABOUT JOINING PROCESSES AND PRINCIPLES. WELDING, BRAZING, AND SOLDERING ARE TRULY INTERDISCIPLINARY ENTERPRISES; NO INDIVIDUAL CAN BE EXPECTED TO BE AN EXPERT IN ALL ASPECTS OF THESE TECHNOLOGIES. THEREFORE, WE HAVE ATTEMPTED TO PROVIDE A HANDBOOK THAT CAN BE USED AS A COMPREHENSIVE REFERENCE BY ANYONE NEEDING MATERIALS-RELATED JOINING INFORMATION. MANY COLLEAGUES AND FRIENDS CONTRIBUTED THEIR TIME AND EXPERTISE TO THIS HANDBOOK, AND WE ARE VERY GRATEFUL FOR THEIR EFFORTS. WE WOULD ALSO LIKE TO EXPRESS OUR THANKS TO THE AMERICAN WELDING SOCIETY FOR THEIR COOPERATION AND ASSISTANCE IN THIS ENDEAVOR. DAVID LEROY OLSON, COLORADO SCHOOL OF MINES THOMAS A. SIEWERT, NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY STEPHEN LIU, COLORADO SCHOOL OF MINES GLEN R. EDWARDS, COLORADO SCHOOL OF MINES OFFICERS AND TRUSTEES OF ASM INTERNATIONAL (1992-1993)

OFFICERS • • • • •

EDWARD H. KOTTCAMP, JR. PRESIDENT AND TRUSTEE SPS TECHNOLOGIES JACK G. SIMON VICE PRESIDENT AND TRUSTEE GENERAL MOTORS CORPORATION WILLIAM P. KOSTER IMMEDIATE PAST PRESIDENT AND TRUSTEE METCUT RESEARCH ASSOCIATES, INC. EDWARD L. LANGER SECRETARY AND MANAGING DIRECTOR ASM INTERNATIONAL LEO G. THOMPSON TREASURER LINDBERG CORPORATION

TRUSTEES • • • • • • • • •

WILLIAM H. ERICKSON FDP ENGINEERING NORMAN A. GJOSTEIN FORD MOTOR COMPANY NICHOLAS C. JESSEN, JR. MARTIN MARIETTA ENERGY SYSTEMS, INC. E. GEORGE KENDALL NORTHROP AIRCRAFT GEORGE KRAUSS COLORADO SCHOOL OF MINES LYLE H. SCHWARTZ NATIONAL INSTITUTE OF STANDARDS & TECHNOLOGY GERNANT E. MAURER SPECIAL METALS CORPORATION ALTON D. ROMIG, JR. SANDIA NATIONAL LABORATORIES MERLE L. THORPE HOBART TAFA TECHNOLOGIES, INC.

MEMBERS OF THE ASM HANDBOOK COMMITTEE (1992-1993) • • • •

ROGER J. AUSTIN (CHAIRMAN 1992-; MEMBER 1984-) CONCEPT SUPPORT AND DEVELOPMENT CORPORATION DAVID V. NEFF (VICE CHAIRMAN 1992-; MEMBER 1986-) METAULLICS SYSTEMS TED L. ANDERSON (1991-) TEXAS A&M UNIVERSITY BRUCE P. BARDES (1993-) MIAMI UNIVERSITY

• • • • • • • • • • • • • • • • • • • • • •

ROBERT J. BARNHURST (1988-) NORANDA TECHNOLOGY CENTRE TONI BRUGGER (1993-) PHOENIX PIPE & TUBE COMPANY STEPHEN J. BURDEN (1989-) CRAIG V. DARRAGH (1989-) THE TIMKEN COMPANY RUSSELL J. DIEFENDORF (1990-) CLEMSON UNIVERSITY AICHA EISHABINI-RIAD (1990-) VIRGINIA POLYTECHNIC & STATE UNIVERSITY GREGORY A. FETT (1993-) DANA CORPORATION MICHELLE M. GAUTHIER (1990-) RAYTHEON COMPANY TONI GROBSTEIN (1990-) NASA LEWIS RESEARCH CENTER SUSAN HOUSH (1990-) DOW CHEMICAL U.S.A. DENNIS D. HUFFMAN (1982-) THE TIMKEN COMPANY S. JIM LBARRA (1991-) AMOCO RESEARCH CENTER J. ERNESTO INDACOCHEA (1987-) UNIVERSITY OF ILLINOIS AT CHICAGO PETER W. LEE (1990-) THE TIMKEN COMPANY WILLIAM L. MANKINS (1989-) INCO ALLOYS INTERNATIONAL, INC. RICHARD E. ROBERTSON (1990-) UNIVERSITY OF MICHIGAN JOGENDER SINGH (1993-) NASA GEORGE C. MARSHALL SPACE FLIGHT CENTER JEREMY C. ST. PIERRE (1990-) HAYES HEAT TREATING CORPORATION EPHRAIM SUHIR (1990-) AT&T BELL LABORATORIES KENNETH TATOR (1991-) KTA-TATOR, INC. MALCOLM THOMAS (1993-) ALLISON GAS TURBINES WILLIAM B. YOUNG (1991-) DANA CORPORATION

PREVIOUS CHAIRMEN OF THE ASM HANDBOOK COMMITTEE • • • • • • • • • • • • • • • • • • • • • • • •

R.S. ARCHER (1940-1942) (MEMBER 1937-1942) L.B. CASE (1931-1933) (MEMBER 1927-1933) T.D. COOPER (1984-1986) (MEMBER 1981-1986) E.O. DIXON (1952-1954) (MEMBER 1947-1955) R.L. DOWDELL (1938-1939) (MEMBER 1935-1939) J.P. GILL (1937) (MEMBER 1934-1937) J.D. GRAHAM (1966-1968) (MEMBER 1961-1970) J.F. HARPER (1923-1926) (MEMBER 1923-1926) C.H. HERTY, JR. (1934-1936) (MEMBER 1930-1936) D.D. HUFFMAN (1986-1990) (MEMBER 1982-1990) J.B. JOHNSON (1948-1951) (MEMBER 1944-1951) L.J. KORB (1983) (MEMBER 1978-1983) R.W.E. LEITER (1962-1963) (MEMBER 1955-1958, 1960-1964) G.V. LUERSSEN (1943-1947) (MEMBER 1942-1947) G.N. MANIAR (1979-1980) (MEMBER 1974-1980) J.L. MCCALL (1982) (MEMBER 1977-1982) W.J. MERTEN (1927-1930) (MEMBER 1923-1933) D.L. OLSON (1990-1992) (MEMBER 1982-1988, 1989-1992) N.E. PROMISEL (1955-1961) (MEMBER 1954-1963) G.J. SHUBAT (1973-1975) (MEMBER 1966-1975) W.A. STADTLER (1969-1972) (MEMBER 1962-1972) R. WARD (1976-1978) (MEMBER 1972-1978) M.G.H. WELLS (1981) (MEMBER 1976-1981) D.J. WRIGHT (1964-1965) (MEMBER 1959-1967)

STAFF ASM INTERNATIONAL STAFF WHO CONTRIBUTED TO THE DEVELOPMENT OF THE VOLUME INCLUDED WILLIAM W. SCOTT, JR., DIRECTOR OF TECHNICAL PUBLICATIONS; SCOTT D.

HENRY, MANAGER OF HANDBOOK DEVELOPMENT; SUZANNE E. HAMPSON, PRODUCTION PROJECT MANAGER; THEODORE B. ZORC, TECHNICAL EDITOR; FAITH REIDENBACH, CHIEF COPY EDITOR; LAURIE A. HARRISON, EDITORIAL ASSISTANT; NANCY M. SOBIE, PRODUCTION ASSISTANT. EDITORIAL ASSISTANCE WAS PROVIDED BY JOSEPH R. DAVIS, KELLY FERJUTZ, NIKKI D. WHEATON, AND MARA S. WOODS. CONVERSION TO ELECTRONIC FILES ASM HANDBOOK, VOLUME 6, WELDING, BRAZING, AND SOLDERING WAS CONVERTED TO ELECTRONIC FILES IN 1998. THE CONVERSION WAS BASED ON THE SECOND PRINTING (1994). NO SUBSTANTIVE CHANGES WERE MADE TO THE CONTENT OF THE VOLUME, BUT SOME MINOR CORRECTIONS AND CLARIFICATIONS WERE MADE AS NEEDED. ASM INTERNATIONAL STAFF WHO CONTRIBUTED TO THE CONVERSION OF THE VOLUME INCLUDED SALLY FAHRENHOLZ-MANN, BONNIE SANDERS, SCOTT HENRY, ROBERT BRADDOCK, AND MARLENE SEUFFERT. THE ELECTRONIC VERSION WAS PREPARED UNDER THE DIRECTION OF WILLIAM W. SCOTT, JR., TECHNICAL DIRECTOR, AND MICHAEL J. DEHAEMER, MANAGING DIRECTOR. COPYRIGHT INFORMATION (FOR PRINT VOLUME) COPYRIGHT © 1993 BY ASM INTERNATIONAL ALL RIGHTS RESERVED. ASM HANDBOOK IS A COLLECTIVE EFFORT INVOLVING THOUSANDS OF TECHNICAL SPECIALISTS. IT BRINGS TOGETHER IN ONE BOOK A WEALTH OF INFORMATION FROM WORLD-WIDE SOURCES TO HELP SCIENTISTS, ENGINEERS, AND TECHNICIANS SOLVE CURRENT AND LONG-RANGE PROBLEMS. GREAT CARE IS TAKEN IN THE COMPILATION AND PRODUCTION OF THIS VOLUME, BUT IT SHOULD BE MADE CLEAR THAT NO WARRANTIES, EXPRESS OR IMPLIED, ARE GIVEN IN CONNECTION WITH THE ACCURACY OR COMPLETENESS OF THIS PUBLICATION, AND NO RESPONSIBILITY CAN BE TAKEN FOR ANY CLAIMS THAT MAY ARISE. NOTHING CONTAINED IN THE ASM HANDBOOK SHALL BE CONSTRUED AS A GRANT OF ANY RIGHT OF MANUFACTURE, SALE, USE, OR REPRODUCTION, IN CONNECTION WITH ANY METHOD, PROCESS, APPARATUS, PRODUCT, COMPOSITION, OR SYSTEM, WHETHER OR NOT COVERED BY LETTERS PATENT, COPYRIGHT, OR TRADEMARK, AND NOTHING CONTAINED IN THE ASM HANDBOOK SHALL BE CONSTRUED AS A DEFENSE AGAINST ANY ALLEGED INFRINGEMENT OF LETTERS PATENT, COPYRIGHT, OR TRADEMARK, OR AS A DEFENSE AGAINST LIABILITY FOR SUCH INFRINGEMENT. COMMENTS, CRITICISMS, AND SUGGESTIONS ARE INVITED, AND SHOULD BE FORWARDED TO ASM INTERNATIONAL. LIBRARY OF CONGRESS CATALOGING-IN-PUBLICATION DATA (FOR PRINT VOLUME) ASM HANDBOOK (REVISED VOL. 6) METALS HANDBOOK. VOLS. 1-2 HAVE TITLE: METALS HANDBOOK. VOL. 4 LACKS ED. STATEMENTS. INCLUDES BIBLIOGRAPHICAL REFERENCES AND INDEXES. CONTENTS: V. 1. PROPERTIES AND SELECTION-IRONS, STEELS, AND HIGH-PERFORMANCE ALLOYS-V. 2. PROPERTIES AND SELECTION-NONFERROUS ALLOYS

AND SPECIAL-PURPOSE MATERIALS-[ETC.]-V. 6. WELDING, BRAZING, AND SOLDERING. 1. METALS-HANDBOOKS, MANUALS, ETC. 2. METAL-WORK-HANDBOOKS, MANUALS, ETC. I. ASM INTERNATIONAL. HANDBOOK COMMITTEE. II. TITLE: METALS HANDBOOK. TA459.M43 1990 620.1'6 90-115 ISBN 0-87170-377-7(V.1) SAN 204-7586 ISBN 0-87170-382-3 PRINTED IN THE UNITED STATES OF AMERICA Energy Sources Used for Fusion Welding Thomas W. Eagar, Massachusetts Institute of Technology

Introduction WELDING AND JOINING processes are essential for the development of virtually every manufactured product. However, these processes often appear to consume greater fractions of the product cost and to create more of the production difficulties than might be expected. There are a number of reasons that explain this situation. First, welding and joining are multifaceted, both in terms of process variations (such as fastening, adhesive bonding, soldering, brazing, arc welding, diffusion bonding, and resistance welding) and in the disciplines needed for problem solving (such as mechanics, materials science, physics, chemistry, and electronics). An engineer with unusually broad and deep training is required to bring these disciplines together and to apply them effectively to a variety of processes. Second, welding or joining difficulties usually occur far into the manufacturing process, where the relative value of scrapped parts is high. Third, a very large percentage of product failures occur at joints because they are usually located at the highest stress points of an assembly and are therefore the weakest parts of that assembly. Careful attention to the joining processes can produce great rewards in manufacturing economy and product reliability. The Section "Fusion Welding Processes" in this Volume provides details about equipment and systems for the major fusion welding processes. The purpose of this Section of the Volume is to discuss the fundamentals of fusion welding processes, with an emphasis on the underlying scientific principles. Because there are many fusion welding processes, one of the greatest difficulties for the manufacturing engineer is to determine which process will produce acceptable properties at the lowest cost. There are no simple answers. Any change in the part geometry, material, value of the end product, or size of the production run, as well as the availability of joining equipment, can influence the choice of joining method. For small lots of complex parts, fastening may be preferable to welding, whereas for long production runs, welds can be stronger and less expensive. The perfect joint is indistinguishable from the material surrounding it. Although some processes, such as diffusion bonding, can achieve results that are very close to this ideal, they are either expensive or restricted to use with just a few materials. There is no universal process that performs adequately on all materials in all geometries. Nevertheless, virtually any material can be joined in some way, although joint properties equal to those of the bulk material cannot always be achieved. The economics of joining a material may limit its usefulness. For example, aluminum is used extensively in aircraft manufacturing and can be joined by using adhesives or fasteners, or by welding. However, none of these processes has proven economical enough to allow the extensive replacement of steel by aluminum in the frames of automobiles. An increased use of composites in aircrafts is limited by an inability to achieve adequate joint strength.

It is essential that the manufacturing engineer work with the designer from the point of product conception to ensure that compatible materials, processes, and properties are selected for the final assembly. Often, the designer leaves the problem of joining the parts to the manufacturing engineer. This can cause an escalation in cost and a decrease in reliability. If the design has been planned carefully and the parts have been produced accurately, the joining process becomes much easier and cheaper, and both the quality and reliability of the product are enhanced. Generally, any two solids will bond if their surfaces are brought into intimate contact. One factor that generally inhibits this contact is surface contamination. Any freshly produced surface exposed to the atmosphere will absorb oxygen, water vapor, carbon dioxide, and hydrocarbons very rapidly. If it is assumed that each molecule that hits the surface will be absorbed, then the time-pressure value to produce a monolayer of contamination is approximately 0.001 Pa · s (10-8 atm · s). For example, at a pressure of 1 Pa (10-5 atm), the contamination time is 10-3 s, whereas at 0.1 MPa (1 atm), it is only 10 × 10-9 s. In fusion welding, intimate interfacial contact is achieved by interposing a liquid of substantially similar composition as the base metal. If the surface contamination is soluble, then it is dissolved in the liquid. If it is insoluble, then it will float away from the liquid-solid interface. Energy Sources Used for Fusion Welding Thomas W. Eagar, Massachusetts Institute of Technology

Energy-Source Intensity One distinguishing feature of all fusion welding processes is the intensity of the heat source used to melt the liquid. Virtually every concentrated heat source has been applied to the welding process. However, many of the characteristics of each type of heat source are determined by its intensity. For example, when considering a planar heat source diffusing into a very thick slab, the surface temperature will be a function of both the surface power density and the time. Figure 1 shows how this temperature will vary on steel with power densities that range from 400 to 8000 W/cm2. At the lower value, it takes 2 min to melt the surface. If that heat source were a point on the flat surface, then the heat flow would be divergent and might not melt the steel. Rather, the solid metal would be able to conduct away the heat as fast as it was being introduced. It is generally found that heat-source power densities of approximately 1000 W/cm2 are necessary to melt most metals.

FIG. 1 TEMPERATURE DISTRIBUTION AFTER A SPECIFIC HEATING TIME IN A THICK STEEL PLATE HEATED

UNIFORMLY ON ONE SURFACE AS A FUNCTION OF APPLIED HEAT INTENSITY; INITIAL TEMPERATURE OF PLATE IS 25 °C (77 °F)

At the other end of the power-density spectrum, heat intensities of 106 or 107 W/cm2 will vaporize most metals within a few microsecond. At levels above these values, all of the solid that interacts with the heat source will be vaporized, and no fusion welding can occur. Thus, the heat sources for all fusion welding processes should have power densities between approximately 0.001 and 1 MW/cm2. This power-density spectrum is shown in Fig. 2, along with the points at which common joining processes are employed.

FIG. 2 SPECTRUM OF PRACTICAL HEAT INTENSITIES USED FOR FUSION WELDING

The fact that power density is inversely related to the interaction time of the heat source on the material is evident in Fig. 1. Because this represents a transient heat conduction problem, one can expect the heat to diffuse into the steel to a depth that increases as the square root of time, that is, from the Einstein equation: X ~ αt

(EQ 1)

where x is the distance that the heat diffuses into the solid, in centimeters: α is the thermal diffusivity of the solid, in cm2/s; and t is the time in seconds. Tables 1 and 2 give the thermal diffusivities of common elements and common alloys, respectively.

TABLE 1 THERMAL DIFFUSIVITIES OF COMMON ELEMENTS FROM 20 TO 100 °C (68 TO 212 °F)

ELEMENT

ALUMINUM ANTIMONY BERYLLIUM BISMUTH CADMIUM CARBON COBALT COPPER GALLIUM GERMANIUM GOLD HAFNIUM

DENSITY

HEAT CAPACITY

THERMAL CONDUCTIVITY

g/cm3

lb/in.3

j/kg · k

calit/g · °c

w/m · k

calit/cm · s · °c

2.699 6.62 1.848 9.80 8.65 2.25 8.85 8.96 5.907 5.323 19.32 13.09

0.098 0.239 0.067 0.354 0.313 0.081 0.320 0.324 0.213 0.192 0.698 0.472

900 205 1880 123 230 691 414 385 331 306 131 147

0.215 0.049 0.45 0.0294 0.055 0.165 0.099 0.092 0.079 0.073 0.0312 0.0351

221 19 147 8 92 24 69 394 29-38 59 297 22

0.53 0.045 0.35 0.020 0.22 0.057 0.165 0.941 0.07-0.09 0.14 0.71 0.053

mm2/s THERMAL DIFFUSIVITY cm2/s 91 0.91 14 0.14 42 0.42 7 0.09 46 0.46 15 0.15 19 0.188 114 1.14 17 0.17 36 0.36 118 1.178 12 0.12

INDIUM IRIDIUM IRON LEAD MAGNESIUM MOLYBDENUM NICKEL NIOBIUM PALLADIUM PLATINUM PLUTONIUM RHODIUM SILICON SILVER SODIUM TANTALUM TIN TITANIUM TUNGSTEN URANIUM VANADIUM ZINC ZIRCONIUM

7.31 22.5 7.87 11.36 1.74 10.22 8.902 8.57 12.02 21.45 19.84 12.44 2.33 10.49 0.9712 16.6 7.2984 4.507 19.3 19.07 6.1 7.133 6.489

0.264 0.813 0.284 0.410 0.063 0.369 0.322 0.310 0.434 0.775 0.717 0.449 0.084 0.379 0.035 0.600 0.264 0.163 0.697 0.689 0.22 0.258 0.234

239 129 460 129 1025 276 440 268 244 131 138 247 678 234 1235 142 226 519 138 117 498 383 280

0.057 0.0307 0.11 0.0309 0.245 0.066 0.105 0.064 0.0584 0.0314 0.033 0.059 0.162 0.0559 0.295 0.034 0.054 0.124 0.033 0.0279 0.119 0.0915 0.067

24 59 75 35 154 142 92 54 70 69 8 88 84 418 134 54 63 22 166 30 31 113 21

0.057 0.14 0.18 0.083 0.367 0.34 0.22 0.129 0.168 0.165 0.020 0.21 0.20 1.0 0.32 0.130 0.150 0.052 0.397 0.071 0.074 0.27 0.050

14 20 21 24 86 50 23.5 23.6 24 24.5 3.0 29 53 170 112 23 38 9 62 13 10 41 12

0.137 0.20 0.208 0.236 0.86 0.50 0.235 0.236 0.24 0.245 0.030 0.286 0.53 1.705 1.12 0.23 0.38 0.092 0.62 0.13 0.10 0.41 0.12

TABLE 2 THERMAL DIFFUSIVITIES OF COMMON ALLOYS FROM 20 TO 100 °C (68 TO 212 °F)

ALLOYS

DENSITY g/cm3

ALUMINUM ALLOYS 1100 2.71 2014 2.80 5052 2.68 6061 2.70 7075 2.80 COPPER ALLOYS COMMERCIAL 8.80 BRONZE CARTRIDGE 8.53 BRASS NAVAL BRASS 8.41 BERYLLIUM 8.23 COPPER 9% ALUMINUM 7.58 BRONZE MAGNESIUM ALLOYS AZ 31 1.78 AZ 91 1.83 ZW 1 1.8 RZ 5 1.84 STAINLESS STEELS

HEAT CAPACITY

THERMAL CONDUCTIVITY

THERMAL DIFFUSIVITY

lb/in.3

j/kg · k

calit/g · °c

w/m · k

calit/cm · s · °c

mm2/s

cm2/s

0.098 0.101 0.097 0.098 0.101

963 963 963 963 963

0.23 0.23 0.23 0.23 0.23

222 193 138 172 121

0.53 0.46 0.33 0.41 0.29

85 71 54 66 45

0.85 0.71 0.54 0.66 0.45

0.318 377

0.09

188

0.45

57

0.57

0.308 377

0.09

121

0.29

38

0.38

0.303 377 0.297 419

0.09 0.1

117 84

0.28 0.20

37 24

0.37 0.24

0.273 435

0.104

60

0.144

18

0.18

0.064 0.066 0.065 0.066

0.25 0.24 0.24 0.23

84 84 134 113

0.20 0.20 0.32 0.27

45 46 74 64

0.45 0.46 0.74 0.64

1050 1005 1005 963

TYPE 301 7.9 TYPE 304 7.9 TYPE 316 8.0 TYPE 410 7.7 TYPE 430 7.7 TYPE 501 7.7 NICKEL-BASE ALLOYS NIMONIC 80A 8.19 INCONEL 600 8.42 MONEL 400 8.83 TITANIUM ALLOYS TI-6AL-4V 4.43 TI-5AL-2.5SN 4.46

0.285 0.285 0.289 0.278 0.278 0.278

502 502 502 460 460 460

0.12 0.12 0.12 0.11 0.11 0.11

16 15.1 15.5 24 26 37

0.039 0.036 0.037 0.057 0.062 0.088

4.1 3.8 3.9 6.7 7.3 10

0.041 0.038 0.039 0.067 0.073 0.10

0.296 460 0.304 460 0.319 419

0.11 0.11 0.10

11 15 22

0.027 0.035 0.052

3.0 3.8 5.8

0.030 0.038 0.058

0.160 611 0.161 460

0.146 0.11

5.9 6.3

0.014 0.015

2.1 3.1

0.021 0.031

For the planar heat source on a steel surface, as represented by Fig. 1, the time in seconds to produce melting on the surface, tm, is given by:

TM = (5000/H.I.)2

(EQ 2)

where H.I. is the net heat intensity (in W/cm2) transferred to the workpiece. Equation 2 provides a rough estimate of the time required to produce melting, and is based upon the thermal diffusivity of steel. Materials with higher thermal diffusivities--or the use of a local point heat source rather than a planar heat source-will increase the time to produce melting by a factor of up to two to five times. On the other hand, thin materials tend to heat more quickly. If the time to melting is considered to be a characteristic interaction time, tI, then the graph shown in Fig. 3 can be generated. Heat sources with power densities that are of the order of 1000 W/cm2, such as oxyacetylene flames or electroslag welding, require interaction times of 25 s with steel, whereas laser and electron beams, at 1 MW/cm2, need interaction times on the order of only 25 μs. If this interaction time is divided into the heat-source diameter, dH, then a maximum travel speed, Vmax, is obtained for the welding process (Fig. 4).

FIG. 3 TYPICAL WELD POOL-HEAT SOURCE INTERACTION TIMES AS FUNCTION OF HEAT-SOURCE INTENSITY.

MATERIALS WITH A HIGH THERMAL DIFFUSIVITY, SUCH AS COPPER OR ALUMINUM, WOULD LIE NEAR THE TOP OF THIS BAND, WHEREAS STEELS, NICKEL ALLOYS, OR TITANIUM WOULD LIE IN THE MIDDLE. URANIUM AND CERAMICS, WITH VERY LOW THERMAL DIFFUSIVITIES, WOULD LIE NEAR THE BOTTOM OF THE BAND.

FIG. 4 MAXIMUM WELD TRAVEL VELOCITY AS A FUNCTION OF HEAT-SOURCE INTENSITY BASED ON TYPICAL HEAT-SOURCE SPOT DIAMETERS

The reason why welders begin their training with the oxyacetylene process should be clear: it is inherently slow and does not require rapid response time in order to control the size of the weld puddle. Greater skill is needed to control the morerapid fluctuations in arc processes. The weld pool created by the high-heat-intensity processes, such as laser-beam and electron-beam welding, cannot be humanly controlled and must therefore be automated. This need to automate leads to increased capital costs. On an approximate basis, the W/cm2 of a process can be substituted with the dollar cost of the capital equipment. With reference to Fig. 2, the cost of oxyacetylene welding equipment is nearly $1000, whereas a fully automated laser-beam or electron-beam system can cost $1 million. Note that the capital cost includes only the energy source, control system, fixturing, and materials handling equipment. It does not include operating maintenance or inspection costs, which can vary widely depending on the specific application. For constant total power, a decrease in the spot size will produce a squared increase in the heat intensity. This is one of the reasons why the spot size decreases with increasing heat intensity (Fig. 4). It is easier to make the spot smaller than it is to increase the power rating of the equipment. In addition, only a small volume of material usually needs to be melted. If the spot size were kept constant and the input power were squared in order to obtain higher densities, then the volume of fused metal would increase dramatically, with no beneficial effect. However, a decreasing spot size, coupled with a decreased interaction time at higher power densities, compounds the problem of controlling the higher-heat-intensity process. A shorter interaction time means that the sensors and controllers necessary for automation must operate at higher frequencies. The smaller spot size means that the positioning of the heat source must be even more precise, that is, on the order of the heat-source diameter, dH. The control frequency must be greater than the travel velocity divided by the diameter of the heat source. For processes that operate near the maximum travel velocity, this is the inverse of the process interaction time, tI (Fig. 3). Thus, not only must the high-heat-intensity processes be automated because of an inherently high travel speed, but the fixturing requirements become greater, and the control systems and sensors must have ever-higher frequency responses. These factors lead to increased costs, which is one reason that the very productive laser-beam and electron-beam welding processes have not found wider use. The approximate productivity of selected welding processes, expressed as length of weld produced per second, to the relative capital cost of equipment is shown in Fig. 5.

FIG. 5 APPROXIMATE RELATIONSHIP BETWEEN CAPITAL COST OF WELDING EQUIPMENT AND SPEED AT WHICH SHEET METAL JOINTS CAN BE PRODUCED

Another important welding process parameter that is related to the power density of the heat source is the width of the heat-affected zone (HAZ). This zone is adjacent to the weld metal and is not melted itself but is structurally changed because of the heat of welding. Using the Einstein equation, the HAZ width can be estimated from the process interaction time and the thermal diffusivity of the material. This is shown in Fig. 6, with one slight modification. At levels above approximately 104 W/cm2, the HAZ width becomes roughly constant. This is due to the fact that the HAZ grows during the heating stage at power densities that are below 104 W/cm2, but at higher power densities it grows during the cooling cycle. Thus, at low power densities, the HAZ width is controlled by the interaction time, whereas at high power densities, it is independent of the heat-source interaction time. In the latter case, the HAZ width grows during the cooling cycle as the heat of fusion is removed from the weld metal, and is proportional to the fusion zone width.

FIG. 6 RANGE OF WELD HAZ WIDTHS AS FUNCTION OF HEAT-SOURCE INTENSITY

The change of slope in Fig. 6 also represents the heat intensity at which the heat utilization efficiency of the process changes. At high heat intensities, nearly all of the heat is used to melt the material and little is wasted in preheating the surroundings. As heat intensity decreases, this efficiency is reduced. For arc welding, as little as half of the heat generated may enter the plate, and only 40% of this heat is used to fuse the metal. For oxyacetylene welding, the heat entering the metal may be 10% or less of the total heat, and the heat necessary to fuse the metal may be less than 2% of the total heat. A final point is that the heat intensity also controls the depth-to-width ratio of the molten pool. This value can vary from 0.1 in low-heat-intensity processes to more than 10 in high-heat-intensity processes. It should now be evident that all fusion welding processes can be characterized generally by heat-source intensity. The properties of any new heat source can be estimated readily from the figures in this article. Nonetheless, it is useful to more fully understand each of the common welding heat sources, such as flames, arcs, laser beams, electron beams, and electrical resistance. These are described in separate articles in the Section "Fusion Welding Processes" in this Volume. Heat Flow in Fusion Welding Chon L. Tsai and Chin M. Tso, The Ohio State University

Introduction DURING FUSION WELDING, the thermal cycles produced by the moving heat source cause physical state changes, metallurgical phase transformation, and transient thermal stress and metal movement. After welding is completed, the finished product may contain physical discontinuities that are due to excessively rapid solidification, or adverse microstructures that are due to inappropriate cooling, or residual stress and distortion that are due to the existence of incompatible plastic strains. In order to analyze these problems, this article presents an analysis of welding heat flow, focusing on the heat flow in the fusion welding process. The primary objective of welding heat flow modeling is to provide a mathematical tool for thermal data analysis, design iterations, or the systematic investigation of the thermal characteristics of any welding parameters. Exact comparisons with experimental measurements may not be feasible, unless some calibration through the experimental verification procedure is conducted. Welding Thermal Process. A physical model of the welding system is shown in Fig. 1. The welding heat source moves at a constant speed along a straight path. The end result, after either initiating or terminating the heat source, is the formation of a transient thermal state in the weldment. At some point after heat-source initiation but before termination, the temperature distribution is stationary, or in thermal equilibrium, with respect to the moving coordinates. The origin of the moving coordinates coincides with the center of the heat source. The intense welding heat melts the metal and forms a molten pool. Some of the heat is conducted into the base metal and some is lost from either the arc column or the metal surface to the environment surrounding the plate. Three metallurgical zones are formed in the plate upon completion of the thermal cycle: the weld-metal (WM) zone, the heated-affected zone (HAZ), and the base-metal (BM) zone. The peak temperature and the subsequent cooling rates determine the HAZ structures, whereas the thermal gradients, the solidification rates, and the cooling rates at the liquid-solid pool boundary determine the solidification structure of the WM zone. The size and flow direction of the pool determines the amount of dilution and weld penetration. The material response in the temperature range near melting temperatures is primarily responsible for the metallurgical changes.

FIG. 1 SCHEMATIC OF THE WELDING THERMAL MODEL

Two thermal states, quasi-stationary and transient, are associated with the welding process. The transient thermal response occurs during the source initiation and termination stages of welding, the latter of which is of greater metallurgical interest. Hot cracking usually begins in the transient zone, because of the nonequilibrium solidification of the base material. A crack that forms in the source-initiation stage may propagate along the weld if the solidification strains sufficiently multiply in the wake of the welding heat source. During source termination, the weld pool solidifies several times faster than the weld metal in the quasi-stationary state. Cracks usually appear in the weld crater and may propagate along the weld. Another dominant transient phenomenon occurs when a short repair weld is made to a weldment. Rapid cooling results in a brittle HAZ structure and either causes cracking problems or creates a site for fatigue-crack initiation. The quasi-stationary thermal state represents a steady thermal response of the weldment in respect to the moving heat source. The majority of the thermal expansion and shrinkage in the base material occurs during the quasi-stationary thermal cycles. Residual stress and weld distortion are the thermal stress and strain that remain in the weldment after completion of the thermal cycle. Relation to Welding Engineering Problems. To model and analyze the thermal process, an understanding of

thermally induced welding problems is important. A simplified modeling scheme, with adequate assumptions for specific problems, is possible for practical applications without using complex mathematical manipulations. The relationship between the thermal behavior of weldments and the metallurgy, control, and distortion associated with welding is summarized below. Welding Metallurgy. As already noted, defective metallurgical structures in the HAZ and cracking in the WM usually

occur under the transient thermal condition. Therefore, a transient thermal model is needed to analyze cracking and embrittlement problems. To evaluate the various welding conditions for process qualification, the quasi-stationary thermal responses of the weld material need to be analyzed. The minimum required amount of welding heat input within the allowable welding speed range must be determined in order to avoid rapid solidification and cooling of the weldment Preheating may be necessary if the proper thermal conditions cannot be obtained under the specified welding procedure. A quasi-stationary thermal model is adequate for this type of analysis. Hot cracking results from the combined effects of strain and metallurgy. The strain effect results from weld-metal displacement at near-melting temperatures, because of solidification shrinkage and weldment restraint. The metallurgical effect relates to the segregation of alloying elements and the formation of the eutectic during the high nonequilibrium solidification process. Using metallurgical theories, it is possible to determine the chemical segregation, the amounts and distributions of the eutectic, the magnitudes and directions of grain growth, and the weld-metal displacement at high temperatures. Using the heating and cooling rates, as well as the retention period predicted by modeling and analysis, hotcracking tendencies can be determined. To analyze these tendencies, it is important to employ a more accurate numerical model that considers finite welding heat distribution, latent heat, and surface heat loss.

Welding Control. In-process welding control has been studied recently. Many of the investigations are aimed at

developing sensing and control hardware. However, a link between weld-pool geometry and weld quality has not been fully established. A transient heat-flow analysis needs to be used to correlate the melted surface, which is considered to be the primary control variable, to the weld thermal response in a time domain. Welding Distortion The temperature history and distortion caused by the welding thermal process creates nonlinear

thermal strains in the weldment. Thermal stresses are induced if any incompatible strains exist in the weld. Plastic strains are formed when the thermal stresses are higher than the material yield stress. Incompatible plastic strains accumulate over the thermal process and result in residual stress and distortion of the final weldment. The material response in the lower temperature range during the cooling cycle is responsible for the residual stresses and weldment distortion. For this type of analysis, the temperature field away from the welding heat source is needed for the modeling of the heating and cooling cycle during and after welding. A quasi-stationary thermal model with a concentrated moving heat source can predict, with reasonable accuracy, the temperature information for the subsequent stress and distortion analysis. Literature Review. Many investigators have analytically, numerically, and experimentally studied welding heat-flow

modeling and analysis (Ref 1, 2, 3, 4, 5, 6, 7, 8, 9, 10, 11, 12, 13, 14, 15, 16, 17, 18). The majority of the studies were concerned with the quasi-stationary thermal state. Lance and Martin (Ref 1), Rosenthal and Schmerber (Ref 2) and Rykalin (Ref 3) independently obtained an analytical temperature solution for the quasi-stationary state using a point or line heat source moving along a straight line on a semi-infinite body. A solution for plates of finite thickness was later obtained by many investigators using the imaged heat source method (Ref 3, 4). Tsai (Ref 5) developed an analytical solution for a model that incorporated a welding heat source with a skewed Gaussian distribution and finite plate thickness. It was later called the "finite source theory" (Ref 6). With the advancement of computer technology and the development of numerical techniques like the finite-difference and finite-element methods, more exact welding thermal models were studied and additional phenomena were considered, including nonlinear thermal properties, finite heat-source distributions, latent heat, and various joint geometries. Tsai (Ref 5), Pavelic (Ref 7), Kou (Ref 8), Kogan (Ref 9), and Brody (Ref 10) studied the simulation of the welding process using the finite-difference scheme. Hibbitt and Marcal (Ref 11), Friedman (Ref 12), and Paley (Ref 13) made some progress in welding simulation using the finite-element method. Analytical solutions for transient welding heat flow in a plate were first studied by Naka (Ref 14), Rykalin (Ref 3), and Masubuchi and Kusuda (Ref 15) in the 1940s and 1950s. A point or line heat source, constant thermal properties, and adiabatic boundary conditions were assumed. Later, Tsai (Ref 16) extended the analytical solution to incorporate Gaussian heat distribution using the principle of superposition. The solution was used to investigate the effect of pulsed conditions on weld-pool formation and solidification without the consideration of latent heat and nonlinear thermal properties. The analysis of the transient thermal behavior of weldments using numerical methods has been the focus of several investigations since 1980. Friedman (Ref 17) discussed the finite-element approach to the general transient thermal analysis of the welding process. Brody (Ref 10) developed a two-dimensional transient heat flow model using a finitedifference scheme and a simulated pulsed-current gas-tungsten arc welding process (GTAW). Tsai and Fan (Ref 18) modeled the two-dimensional transient welding heat flow using a finite-element scheme to study the transient welding thermal behavior of the weldment. General Approach. The various modeling and analysis schemes summarized above can be used to investigate the

thermal process of different welding applications. With adequate assumptions, analytical solutions for the simplified model can be used to analyze welding problems that show a linear response to the heat source if the solutions are properly calibrated by experimental tests. Numerical solutions that incorporate nonlinear thermal characteristics of weldments are usually required for investigating the weld-pool growth or solidification behavior. Numerical solutions can also be necessary for metallurgical studies in the weld HAZ if the rapid cooling phenomenon is significant under an adverse welding environment, such as welding under water. Thermally related welding problems can be categorized as: • • •

SOLIDIFICATION RATES IN THE WELD POOL COOLING RATES IN THE HAZ AND ITS VICINITY THERMAL STRAINS IN THE GENERAL DOMAIN OF THE WELDMENT

The domain of concern in the weld pool solidification is within the molten pool area, in which the arc (or other heat source) phenomena and the liquid stirring effect are significant. A convective heat-transfer model with a moving boundary at the melting temperature is needed to study the first category, and numerical schemes are usually required, as well. The HAZ is always bounded on one side by the liquid-solid interface during welding. This inner-boundary condition is the solidus temperature of the material. The liquid weld pool might be eliminated from thermal modeling if the interface could be identified. A conduction heat-transfer model would be sufficient for the analysis of the HAZ. Numerical methods are often employed and very accurate results can be obtained. The thermal strains caused by welding thermal cycles are caused by the nonlinear temperature distribution in the general domain of the weldment. Because the temperature in the material near the welding heat source is high, very little stress can be accumulated from the thermal strains. This is due to low rigidity, that is, small modulus of elasticity and low yield strength. The domain for thermal strain study is less sensitive to the arc and fluid-flow phenomena and needs only a relatively simple thermal model. Analytical solutions with minor manipulations often provide satisfactory results. In this article, only the analytical heat-flow solutions and their practical applications are addressed. The numerical conduction solutions and the convective models for fluid flow in a molten weld pool are not presented.

References

1. N.S. BOULTON AND H.E. LANCE-MARTIN, RESIDUAL STRESSES IN ARC WELDED PLATES, PROC. INST. MECH. ENG., VOL 33, 1986, P 295 2. D. ROSENTHAL AND R. SCHMERBER, THERMAL STUDY OF ARC WELDING, WELD. J., VOL 17 (NO. 4), 1983, P 2S 3. N.N. RYKALIN, "CALCULATIONS OF THERMAL PROCESSES IN WELDING," MASHGIZ, MOSCOW, 1951 4. K. MASUBUCHI, ANALYSIS OF WELDED STRUCTURES, PERGAMON PRESS, 1980 5. C.L. TSAI, "PARAMETRIC STUDY ON COOLING PHENOMENA IN UNDERWATER WELDING," PH.D THESIS, MIT, 1977 6. C.L. TSAI, FINITE SOURCE THEORY, MODELING OF CASTING AND WELDING PROCESSES II, ENGINEERING FOUNDATION MEETING, NEW ENGLAND COLLEGE (HENNIKER, NH), 31 JULY TO 5 AUG 1983, P 329 7. R. PAVELIC, R. TANAKUCHI, O. CZEHARA, AND P. MYERS, EXPERIMENTAL AND COMPUTED TEMPERATURE HISTORIES IN GAS TUNGSTEN ARC WELDING IN THIN PLATES, WELD. J., VOL 48 (NO. 7), 1969, P 295S 8. S. KOU, 3-DIMENSIONAL HEAT FLOW DURING FUSION WELDING, PROC. OF METALLURGICAL SOCIETY OF AIME, AUG 1980, P 129-138 9. P.G. KOGAN, THE TEMPERATURE FIELD IN THE WELD ZONE, AVE. SVARKA, VOL 4 (NO. 9), 1979, P 8 10. G.M. ECER, H.D. DOWNS, H.D. BRODY, AND M.A. GOKHALE, HEAT FLOW SIMULATION OF PULSED CURRENT GAS TUNGSTEN ARC WELDING, MODELING OF CASTING AND WELDING PROCESSES, ENGINEERING FOUNDATION 1980 MEETING (RINDGE, NH), 3-8 AUG 1980, P 139160 11. H. HIBBITT AND P. MARCAL, A NUMERICAL THERMOMECHANICAL MODEL FOR WELDING AND SUBSEQUENT LOADING OF A FABRICATED STRUCTURE, COMPUT. STRUCT., VOL 3, 1973, P 1145 12. E. FRIEDMAN, THERMOMECHANICAL ANALYSIS OF THE WELDING PROCESS USING FINITE ELEMENT METHODS, TRANS. ASME, AUG 1975, P 206 13. Z. PALEY AND P. HIBBERT, COMPUTATION OF TEMPERATURE IN ACTUAL WELD DESIGN, WELD. J., VOL 54 (NO. 11), 1975, P 385.S

14. T. NAKA, TEMPERATURE DISTRIBUTION DURING WELDING, J. JPN. WELD. SOC., VOL 11 (NO. 1), 1941, P 4 15. K. MASUBUCHI AND T. KUSUDA, TEMPERATURE DISTRIBUTION OF WELDED PLATES, J. JPN WELD. SOC., VOL 22 (NO. 5), 1953, P 14 16. C.L. TSAI AND C.A. HOU, THEORETICAL ANALYSIS OF WELD POLL BEHAVIOR IN THE PULSED CURRENT GTAW PROCESS, TRANSPORT PHENOMENA IN MATERIALS PROCESSING, ASME WINTER ANNUAL MEETING, 1983 17. E. FRIEDMAN, "FINITE ELEMENT ANALYSIS OF ARC WELDING," REPORT WAPD-TM-1438, DEPARTMENT OF ENERGY, 1980 18. J.S. FAN AND C.L. TSAI, "FINITE ELEMENT ANALYSIS OF WELDING THERMAL BEHAVIOR IN TRANSIENT CONDITIONS," 84-HT-80, ASME Heat Flow in Fusion Welding Chon L. Tsai and Chin M. Tso, The Ohio State University

Mathematical Formulations Conduction Equation. A diagram of the welding thermal model is shown in Fig. 1. The origin of the moving coordinates (w,x,z) is fixed at the center of the welding heat source. The coordinates move with the source at the same speed. The conduction equation for heat flow in the weldments is:

∇.(λ∇θ ) + ρ C p v

∂θ • ∂θ + Q = ρC pv ∂w ∂t

(EQ 1)

The initial condition is:

θ= θ 0, AT T = 0

(EQ 2)

and the general boundary condition is: λ

• ∂θ ∂θ ∂θ Iw + Iy + I z − q + h(θ − θ ∞ ) = 0 ∂w ∂y ∂z

(EQ 3)

where ∇ is a differential operator; θ is the temperature; θ ∞ is the environmental temperature; θ0 is the initial temperature; λ is thermal conductivity; ρ is density; Cp is specific heat h is the surface heat-loss coefficient; lw, ly, and lz are the direction cosines of the boundary surface; Q is the volumetric heat source, t is time, and v is welding speed. The volumetric heat source represents the Joule heating in the weldment that is due to the electric current flow within that conducting medium. The total energy of such heating in welding is usually minimal, compared to the arc heat input. The majority of the energy is concentrated in a very small volume beneath the arc (Ref 5). In other words, a very high energy density generation exists in the weld pool, and it may have a significant effect on transient pool growth and solidification. Heat-Source Formulation. The direction cosines on the surface that receive the heat flux from the welding source (z =

0) are lw = ly = 0 and lz = -1. Within the significant heat-input area (to be defined later in this section), the heat loss coefficient, h, is zero. The distribution of the welding heat flux on the weldment surface can be characterized, in a general form, by a skewed Gaussian function (Ref 19): • • βv   q (r , w) = q 0 exp  −Cr ² − w 2k  

(EQ 4)

•

where β is a weight constant, κ is the thermal diffusivity of the base material, C is a shape constant, q is heat flux as a •

function of (r, w), q 0 is heat flux at the source center, r is the radial coordinate from the source center, and v is welding speed. The weight constant, β, indicates the significance of the welding travel speed. A normal distribution of the welding heat flux is obtained if the weight constant is zero. In general, the total energy input to the weldment, which is a fraction of the total welding power generated by the welding machine, is the sum of the concentrated heat and the diffused heat (Ref 20). The concentrated heat is carried by the core of the energy transmission medium, for example, the arc plasma column. The diffused heat reaches the weld surface by radiation and convection energy transport from the core surface. The heat-flux distribution is a function of the proportional values between these two types of energy. The fraction of the total welding power reaching the weldment indicates the heating efficiency of the welding process, and the fraction percentage is defined as welding heat efficiency, η. The shape constant, C, can be obtained in terms of the core diameter, D, and the concentration factor, F. The concentration factor is defined as the ratio of the concentrated heat to the net energy reaching the weldment. The core diameter can be assumed to be the diameter of the plasma column in the arc welding process. The concentration factor and welding heat efficiency are not fully understood and have been subjected to manipulation during the mathematical analyses in order to obtain a better correlation with the experimental data. Assuming a normal heat-flux model, two concentrations are required to determine the shape constant and the heat flux at •

the source center, q 0. By integrating Eq 4 over the core heat area and the entire heat input domain (r = 0 → ∞ ), the shape factor can be determined by dividing the two integrals. The heat flux at the source center can then be determined from the second integral. The two constants are expressed as: C=

4larc [1/(1 − F )] D²

(EQ 5)

where larc is the length of the arc plasma column. •

q0 •

=

Q

C π

(EQ 6)

In the case of arc welding, •

q0 C = η EI π

(EQ 7)

where E is the welding arc voltage and I is the welding current. For practical purposes, the welding heat source can be considered to be restricted within a circle of radius ra, where the •

heat flux drops to 1/100 of the center flux q 0. The radius of the significant heat input area can be written as:  I 100  ra =  arc   C 

0.5

(EQ 8)

Surface Heat Loss. The heat-loss coefficient, h, represents both radiation and convection heat loss from the boundary surfaces outside the significant heat input area. The formulation for both heat-loss mechanisms can be written as the radiation heat-loss coefficient (in air):

HRAD = εσ(θw +

∞ )(

θW 2- θ ∞ 2))

(EQ 9)

or the natural convection heat-loss coefficient (in air): H AC = 0.00042

θW − θ ∞ B

(EQ 10)

or the convection heat-loss coefficient (in water):

HCW = 0.442(θW - θ ∞ )0.25

(EQ 11)

where ε is emissivity, σ is the Stefan-Boltzmann constant, θw is the surface temperature, θ ∞ , is the environmental temperature, and B is the characteristic surface dimension. Natural convection is dominant at a temperature below 550 °C (1020 °F), whereas radiation becomes more important at temperatures above this level. The total heat-loss coefficient is the sum of Eq 9 and 10. The characteristic surface dimension is the effective distance from the source beyond which the temperature rises insignificantly during welding. The characteristic dimension for steel is about 150 mm (6 in.) (Ref 5). In underwater welding, heat losses are primarily due to heat transfer from the surface to the moving water environment. This motion is created by the rising gas column in the arc area (Ref 21). For an insulated surface, no heat transfer into or out of the surface is assumed. The temperature gradient normal to the surface is zero, and can be represented by:

N· ∇θ=0

(EQ 12)

where n is a unit vector normal to the surface and equals (lw2 + ly2 + lz2)0.5. Other Boundary Conditions. There are several other possible boundary conditions in welding heat-flow modeling that

depend on the assumptions used for model simplification. One is the condition at infinity: θ =θ ∞ or lim r →∞

∂θ =0 ∂r

(EQ 13)

Another is the condition near the heat source. In the case of a line source for a thin plate: -2πλ H lim r r →∞

∂θ = η EI ∂r

(EQ 14)

In the case of a point source for a thick plate: -2πλ lim r ² r →∞

∂θ = η EI ∂r

(EQ 15)

In the case of a finite source for a thick plate: •

N · (-λ ∇ θ) = q ; R ≤ RA Another is represented by the conditions at the solid-liquid interface:

(EQ 16)

θ1 = θS = θM n.(λ∇θ )s − n(λ∇θ )1 = ± ρ s L

(EQ 17) ds dt

(EQ 18)

where + indicates the melting process and - indicates the solidification process. The subscripts s and l indicate the temperature and the properties in a solid and liquid, respectively. The n is a normal vector on the boundary surface or interface, ra is the radius of the heat-input area, L is the latent heat of the base material, and the subscript m represents the melting temperature of the base material.

References cited in this section

5. C.L. TSAI, "PARAMETRIC STUDY ON COOLING PHENOMENA IN UNDERWATER WELDING," PH.D THESIS, MIT, 1977 19. R.L. APPS AND D.R. MILNER, HEAT FLOW IN ARGON-ARC WELDING, BR. WELD. J., VOL 2 (NO. 10), 1955, P 475 20. H.S. CARSLAW AND J.C. JAEGER, CONDUCTION OF HEAT IN SOLIDS, OXFORD PRESS 21. C.L. TSAI AND J.H. WU, "AN INVESTIGATION OF HEAT TRANSPORT PHENOMENA IN UNDERWATER WELDING," PRESENTED AT THE ASME WINTER ANNUAL MEETING (MIAMI BEACH, FL), 1985

Heat Flow in Fusion Welding Chon L. Tsai and Chin M. Tso, The Ohio State University

Engineering Solutions and Empirical Correlation General Solutions. The general (analytical) heat-flow solutions for fusion welding can be categorized by those appropriate for a thick plate, a thin plate, or a plate with finite thickness. In most cases, the boundary surfaces (except for the heat-input area) are assumed to be adiabatic, and the thermal properties are independent of temperature. The various metallurgical zones in the weldment are assumed to be homogeneous, and the thermal model is linear.

The solutions give the temperature for a specific point if the welding velocity, v, voltage, E, and current, I, as well as the physical properties of the plate material (ρ, λ, Cp) and the welding heat efficiency, η, are known. This specific point is defined by r and w in: r = w² + y ² + z ²

(EQ 19)

where w = x - vt. The heat-flow solutions are not accurate at points near the welding arc, because a point source or line source is assumed for thick and thin plates, respectively. To approximate the transient temperature changes at the start and end of a weld, Fig. 2 shows a global coordinate system (x,y,z), the origin of which is fixed at the source initiation, where t0 is the welding time and t1 is the time after the welding heat-source termination. The temperature solutions at t0 and t1 are the temperature changes at the start and end of the weld, respectively.

FIG. 2 GLOBAL AND MOVING COORDINATE SYSTEMS FOR WELDING HEAT CONDUCTION.

The temperature solution for thick plate at the arc start location is: η EI 2πλ vto

θ − θ0 =

(EQ 20)

The quasi-stationary temperature distribution is: θ − θ0 =

η EI  −v ( w + r )  exp   2πλ r  2κ 

(EQ 21)

At the arc termination location, the solution is: θ − θ0 =

η EI 2πλ vt1

(EQ 22)

The temperature solution for thin plate at the arc start location is:

θ − θ0 =

v ²t v ²t η EI exp o ko o 2πλ H 2κ 2κ

(EQ 23)

The quasi-stationary temperature distribution is: θ − θ0 =

η EI −vw vr exp ko 2πλ H 2κ 2κ

(EQ 24)

For the arc termination location, the solution is: θ − θ0 =

v ²t η EI  v ²t  exp 1 ko  1  2πλ H 2κ  2κ 

(EQ 25)

where K0 is the modified Bessel function of the second kind of zeroth order and η EI is the welding heat input rate. Temperature for Plate With Finite Thickness. The image method enables the investigator to superimpose the

solutions for an infinitely thick plate, the source of which is placed on imaginary surfaces until the proper boundary

conditions on the plate surfaces are obtained. This method is based on the premise that if a solution satisfies the governing equation and the boundary conditions, then it must be not only a correct solution, but the only solution (that is, the uniqueness of solution premise). Using the image method, the solution for plates of finite thickness with a adiabatic surfaces can be modified from the respective temperature solutions described previously. Let φ 0 (w,y,z,t) be the initial solution for an infinitely thick plate. The temperature solution for a finite thick plate can be obtained by super imposing the imaginary solutions φ mn (w,ym, zn, t) and φ 'mn (w,y' m, z'n, t) to the initial solution, and this can be written in a general form as: ∞

∞

θ − θ 0 = −φ0 ( w, y, z , f ) + ∑ ∑ [φmn ( w, ym , zm , t ] + φ 'mn ( w, y 'm , Z m t )]

(EQ 26)

m=0 n=0

where ym = 2mB - y; y'm = 2mB + y; zn = 2nH - z; and z'n = 2nH + z, in which B is the half width and H is the thickness of the plate. The subscripts m and n are integers that vary from zero to infinity. For a plate with sufficient width, the subscript m is zero. The solution will coverage and reach the correct adiabatic surface condition in six to ten superposition steps, depending on the thickness of the plate. The two-dimensional solution (that is, thin plate) is generally used for any solution that requires more than ten superposition steps. Equation 26 can be expressed as:   −vr2 n  −vw    −vr  η EI exp    exp   ∞  exp  2k 2 k 2 k     +   θ − θ0 = −  ∑ 2πλ r r2 n  n = 1    

   −vr2 n +1     exp  2k    +    r2 n −1     

(EQ 27)

where

R2N = W2 + Y2 + (2NH - Z)2 AND R2N+1 = W2 + Y2 + (2NH + Z)2

(EQ 28)

Cooling Rate. Frequently, it is desirable to know the cooling rate experienced at some location in a weldment to enable a

prediction of the metallurgical structure in that area. A general methodology by which cooling rate equations are obtained from the temperature-distribution equations is discussed below. Recall that the moving coordinate w is defined by w = x - vt. Using this definition, it is easily shown that: ∂w = −v ∂t

(EQ 29)

Using the chain rule, the cooling-rate equation is: ∂θ ∂θ = −v ∂t ∂w

(EQ 30)

Because the temperature-distribution equations are a function of w and r, the cooling-rate equations can be obtained by differentiating the temperature-distribution equations with respect to w and multiplying by -v. The cooling rate is defined as the slope of a tangent line drawn on the temperature-time curve. Because the cooling rate changes with temperature, when one speaks of a cooling rate, the specific temperature, θc, at which it occurs must also be

given. In a weldment, the variable of interest is the cooling rate at the critical temperature that ultimately defines what type of metallurgical structure will result (if the material is heat treatable). For steels, this critical temperature is the "nose" of the continuous cooling-transformation (CCT) curve. At this temperature, the cooling rate determines if upper transformation products (pearlite, upper bainite) or lower transformation products (martensite, lower bainite) will form. For many steels, this critical temperature ranges from approximately 200 to 540 °C (400 to 1000 °F). The cooling rate in a weldment is also a function of location. In order to find a cooling-rate equation, the particular location in the weldment that is of interest must be defined. The resulting cooling-rate equation will be applicable only to that location. The differentiation, ∂ θ/ ∂ w, of either Eq 21 or 27, which is required to obtain the cooling-rate expression, will result in a function of w and r. The variable r can be written in terms of w if the location of interest is defined by a given set of values of y and z. This relationship for r, once formulated, can then be substituted into ∂ θ/ ∂ w, the result being a function of w alone. To determine w corresponding to the critical temperature, θc, a temperature-distribution equation is required (Eq 21 or 27). The aforementioned r-w relationship and temperature distribution equation (Eq 21 or 27) where θ is equal to θc, critical temperature, are used to determine w. Then w is substituted into the dθ/dw expression obtained previously. The end result will be an equation that defines the cooling rate for a particular location in the weldment, and, being a function of the critical temperature, the welding conditions and thermal conductivity of the base material. To determine the cooling rate in a thick plate along the weld centerline (that is y = 0) for a particular critical temperature, the cooling-rate equation can be reduced to: ∂θ 2πλ v (θ c − θ 0 ) = ∂t η EI

2

(EQ 31) •

This equation has been used to predict weld cooling rates in shop practices. Cooling rate is inversely (that is, Q /V) and is proportional to thermal conductivity and the critical temperature at which the cooling rate needs to be evaluated. On the basis of experimental results, a cooling-rate equation was developed for the HAZ of low-carbon steel weldments (Ref 25). This equation considers the combined effects of plate thickness, H, preheating temperature, θ0, and welding conditions, and is given as: ∂θ ∂t

HAZ

 θ − θ 0 1.7  2 −1  H − H o = 0.35  c 1 + tan    α  I / v   π

       

0.8

(EQ 32)

The variables α and H0 depend on the critical temperature of interest. Several values are given in Table 1.

TABLE 1 SELECTED CRITICAL TEMPERATURE AND CORRESPONDING VALUES FOR α AND H0

CRITICAL TEMPERATURE, θC °C °F 700 1290 540 1000 300 570

α

H0 mm 9.9 14.2 19.8

in. 0.39 0.56 0.78

mm 2.0 4.1 9.9

in. 0.08 0.16 0.39

The units used in Eq 32 are important, because the same units that were used in developing the equation must be employed in its application. The plate thickness, H, must be given in inches, and the travel speed, v, must be given in in./min. The temperatures θc and θ0 must be given in °C, and the welding current, I, must be given in amperes. Using the

correct units, the application of Eq 32 will result in a predicted cooling rate (°C/s) for the HAZ of a low-carbon steel weldment. For low-carbon steels welded by the shielded metal arc welding (SMAW), gas-metal arc welding (GMAW), and submerged arc welding (SAW) processes, an empirical equation has been developed that correlates the weld-metal cooling rate at 538 °C (1000 °F) with a 95 to 150 °C (200 to 300 °F) preheat, the weld nugget area (Ref 26): 1.119

  2012 °C / s =    nugget area 

(EQ 33)

where the nugget area is in mm2. For low-carbon steels, an empirical chart for determining nugget area for a given welding condition (Fig. 3) also has been developed (Ref 27). The straight line drawn to connect the current and travel speed intersects the nugget area at the calculated value. Welding voltage controls weld bead shape.

FIG. 3 RELATION BETWEEN NUGGET AREA, HEAT INPUT, AND CURRENT

Peak Temperature. An equation to determine the peak-temperature in a weldment at a given distance y from the weld

centerline would enable the prediction of HAZ sizes, as well as weld bead widths. The general concept of obtaining a peak-temperature equation, as well as some results that have been obtained, are discussed below. Consider Fig. 4 and note that the maximum, or peak, temperature is given when ∂ θ/ ∂ t = 0. For the thick-plate model, the cooling rate can be obtained by differentiating Eq 21 and multiplying by -v: ∂θ ∂θ −vη EI  −v ( w + r )   − w v  w   = −v = exp   x  r ² − 2k  1 + r   ∂t ∂w 2πλ r  2k    

Clearly, the only way that ∂ θ/ ∂ t can be equal to zero is if

(EQ 34)

−w v  w  − 1 + = 0  r ² 2k  r 

(EQ 35)

FIG. 4 SCHEMATIC SHOWING PEAK TEMPERATURE AT (WP, YP, ZP) WITH WP TO BE DETERMINED FOR A GIVEN PEAK TEMPERATURE VALUE FOR A GIVEN (YP, ZP) LOCATION. (A) ISOTHERMS. (B) TEMPERATURE HISTORY

Equation 35 describes the relationship that must exist between the two location variables, r and w, for the temperature at the point to be equal to the peak temperature. If this expression were to be substituted into the temperature distribution equation for thick plates (Eq 21) and solved for w and r (two equations and two unknowns), then the location of the peak temperature could be determined in terms of w and r. The location given by r and w would be easily converted to y and z as:

R2 = W2 + Y2 + Z2

(EQ 36)

Such a solution for r and w is not explicitly possible, however, because the equations for r and w that result are not explicit. Consequently, iterative techniques are required, resulting in a solution that is both cumbersome and timeconsuming. One method of obtaining a simpler thick-plate peak-temperature equation is to assume that the heat input is from an instantaneous line on the surface of the plate, rather than from a moving point source (that is, v → ∞ ). This allows the elimination of the time dependency in the peak-temperature evaluation. Using this assumption, the temperature distribution is given by: θ − θ0 =

η EI  r²  exp   2πλ t  4κ t 

(EQ 37)

Again, to find the peak-temperature location, ∂ θ/ ∂ t is set equal to zero and the equation is solved for r. The result is: r² =1 4κ t

(EQ 38)

Substituting Eq 38 into Eq 37 yields the peak-temperature expression:  e ρ C pπ r ²   η EI  1 = /  θ − θ0  2   v 

(EQ 39)

It has been found that Eq 39 gives results that are too high, but that the slope of 1/(θp -θ0) versus r2 is accurate. To rectify this situation, Eq 39 is forced to fit experimental results by specifying a known temperature/location condition. When this is done, Eq 39 becomes:    [e ρ C π (r ² − r 2 )] / 2  1 1 p t + = EI η θ p − θ0     θr − θ0      v   

(EQ 40)

where θr and rr are the reference temperature and distance. If the peak temperature (θp) evaluation is restricted to locations on the plate surface (z = 0), and if the reference temperature and distance are assumed to be the melting temperature and the distance from the weld centerline to the fusion boundary (one half of the weld bead width), then Eq 40 can be written as: 2  d   e ρ C pπ ( y ² −   2  1  2 = θ p − θ0   η EI      v    

   1   + θ −θ  m 0   

(EQ 41)

where θm is the melting temperature and d is the weld bead width. This equation gives the peak temperature θp in a thick plate at a distance y from the weld centerline. Solidification Rate. The weld solidification structure can be determined by using the constitutional supercooling

criterion. Three thermal parameters that influence the solidification structure are temperature gradient normal to the solidliquid interface, G (°C/cm), solidification rate of the interface, R (cm/s), and cooling rate at the interface, dθ/dt at melting temperature (°C/s and equal to the product of GR). The microstructure may change from being planar to being cellular, a columnar dendrite, or an equiaxial structure if the G/R ratio becomes smaller. The dendrite arm spacing will decrease as the cooling rate increases. The solidification structure becomes refined at higher cooling rates.

At a quasi-steady state, the weld pool solidifies at a rate that is equal to the component of the electrode travel speed normal to the solid-liquid interface. Therefore, the solidification rate varies along the solid-liquid interface from the electrode travel speed, at the weld trailing edge, to zero, at the maximum pool width. The temperature gradient and the cooling rate at the solid-liquid interface can be determined from Eq 21, 24, and 27. Modified Temperature Solution. The temperature solutions have a singularity at the center of the heat source. This singularity causes the predicted temperatures to be inaccurate in the area surrounding the heat source. However, a condition exists in which the peak temperature along the weld bead edge, that is, the solid-liquid interface location at the maximum pool width, is the melting temperature of the material. Using this temperature condition as a boundary condition for the temperature solutions, Eq 21 and 24 can be modified as shown below for thin plate:  −vw   −vr  θ − θ 0 = Bz exp   K0    2κ   2κ 

(EQ 42)

where Bz is a heat input constant to be determined from the weld bead width, d.  K (vr / 2κ )  θ m − θ 0 exp(vrB / 2κ )  0 B   K1 (vrB / 2κ )  Bz = K 0 (vrB / 2κ )

(EQ 43)

with a bead width of:   K (vr / 2κ )  2  d = 2rB 1 −  0 B   K (vr / 2κ )     1 B 

0.5

(EQ 44)

and for thick plate:  −v ( w + r )  θ m − θ 0 = Bz exp  /r  2κ 

(EQ 45)

 1  −vrB   − 1 Bz = (θ m − θ 0 )rB exp     2κ  1 + (2κ / vrB ) 

(EQ 46)

with a bead width of:   1 d = 2rB 1 −   1 + (2κ / vrB ) 

0.5

(EQ 47)

The welding heat input, Q, is replaced by the weld bead width. A Practical Application of Heat Flow Equations (Ref 22). The thermal condition in and near the weld metal must be

established to control the metallurgical events in welding. The particular items of interests are the: • • •

DISTRIBUTION OF PEAK TEMPERATURE IN THE HAZ COOLING RATES IN THE WELD METAL AND IN THE HAZ SOLIDIFICATION RATE OF THE WELD METAL

Although the following discussion primarily focuses on manual are welding, certain general statements are applicable to all welding processes.

Peak Temperatures. The distribution of peak temperatures in the base metal adjacent to the weld is given by (Ref 23): 4.13ρ C p tY 1 1 = + Tp − T0 H met Tm − T0

(EQ 48)

where Tp is the peak temperature (°C) at distance Y (mm) from the weld fusion boundary, T0 is the initial temperature (°C), Tm is the melting temperature (°C), Hnet is the net energy input equal to ηEI/v (J/s · mm), ρ is the density of the material (g/mm3), Cp is the specific heat of solid metal (J/g · °C), and t is the thickness of the base metal (mm). Equation 48 can be used in order to determine the: • • •

PEAK TEMPERATURES AT SPECIFIC LOCATIONS IN THE HAZ WIDTH OF THE HAZ EFFECT OF PREHEAT ON THE WIDTH OF THE HAZ

In addition, determination of the peak temperature at specific locations in the HAZ and the width of the HAZ can be obtained by the procedure described from Eq 34, 35, 36, 37, 38, 39, 40, and 41. Cooling Rate. Because the cooling rate varies with position and time, its calculation requires the careful specification of

conditions. The most useful method is to determine the cooling rate on the weld centerline at the instant when the metal passes through a particular temperature of interest, Tc. At a temperature well below melting, the cooling rate in the weld and in its immediate HAZ is substantially independent of position. For carbon and low-alloy steels, Tc is the temperature near the pearlite "nose" temperature on the time-temperature transformation (TTT) diagram. The value of Tc = 550 °C (1020 °F) is satisfactory for most steels, although not critical. The cooling rate for thick plate (Ref 23) is: R=

2πλ (Tc − T0 )² H met

(EQ 49)

Equation 49 is comparable to Eq 31, which was obtained by the procedure described in the section "Cooling Rate" in this article. Adams (Ref 24) has developed a cooling rate equation for thin plate along the centerline from Eq 24. The cooling rate for thin plate (Ref 24) is: 2

 t  R = 2πλρ C p   + (Tc − T0 )³  H met 

(EQ 50)

where R is the cooling rate (°C/s) at a point on the weld centerline at just that moment when the point is cooling past the Tc, and λ is the thermal conductivity of the metal (J/mm · s · °C). The dimensionless quantity τ , called the "relative plate thickness," can be used to determine whether the plate is thick or thin:

τ =t

ρ C p (Tc − T0 ) H met

(EQ 51)

The thick-plate equation applies when τ is greater than 0.75, and the thin-plate equation applies when τ is less than that value.

Equations 49 and 50 are used to determine the cooling rate along the centerline for thick plate and thin plate, respectively. If one is interested in the cooling rate at the location at distance y (in mm) from the centerline, iterative techniques should be used to solve the cooling rate. First, w and r can be obtained by iteration of the simultaneous equation, which consists of Eq 21 or 27, where θ equals θc and r2 = w2 + y2, where y is given. Then substitute w and r into the differentiation, ∂ θ/ ∂ t = -v ∂ θ/ ∂ w, from the temperature from Eq 21 or 27. The result will be the cooling rate for thick or thin plate located at y distance from the centerline: ∂θ ∂t θ c

In addition, the cooling rate for HAZ of low-carbon steel weldments can be obtained from Eq 32 directly. The solidification rate can have a significant effect on metallurgical structure, properties, response to heat treatment, and soundness. The solidification time, St, of weld metal, measured in seconds, is: St =

LH net 2πλ C p (Tm − To )²

(EQ 52)

where L is the heat of fusion (J/mm3). Example 1: Welding of 5 mm (0.2 in.) Thick Low-Carbon Steels. The thermal properties needed for heat flow analysis are assumed to be:

MELTING TEMPERATURE (TM), °C (°F) AUSTENIZATION TEMPERATURE, °C (°F) THERMAL CONDUCTIVITY (λ), W/M · K (J/MM · S · °C) VOLUMETRIC SPECIFIC HEAT (ρCP), J/MM3 · °C HEAT OF FUSION (L), J/MM3

1510 (2750) 730 (1350) 11.7 (0.028) 0.0044 2

The welding condition is assumed to be:

CURRENT (I), A ARC VOLTAGE (E), V TRAVEL SPEED (V), MM/S (IN./S) PREHEAT (T0), °C (°F) HEAT-TRANSFER EFFICIENCY (η) NET ENERGY INPUT, HNET, J/MM (KJ/IN.)

200 20 5 (0.2) 25 (77) 0.9 720 (18.3)

Calculation of the HAZ Width. The value of Y at Tp = 730 °C (1345 °F) must be determined from Eq 48: 4.13(0.0044)5Yz 1 1 = 730 − 25 720 1510 − 25

resulting in a value for Yz (the width of the HAZ) of 5.9 mm (0.24 in.).

(EQ 53)

In addition, the HAZ width can be obtained by the procedure described in Eq 34, 35, 36, 37, 38, 39, 40, and 41. By substituting w = wp, r = rp, y = yp, and z = 0 (on surface) into Eq 35 and 36, we obtain:

− wp rp2

−

v  wp 1 + 2κ  rp

  = 0 

rp2 = w2p + y 2p

(EQ 54) (EQ 55)

The substitution of Eq 55 in Eq 54 yields:

− wp w +y 2 p

2 p

=

wp v 1+ 2κ w2p + y 2p

(EQ 56)

Inserting v = 5 mm/s, κ= λ/ρcp = 0.028/0.004 into Eq 56:

− wp w +y 2 p

2 p

= 0.3831 +

wp w2p + y 2p

(EQ 57)

wp and yp should satisfy Eq 57. From the temperature distribution equation (Eq 21) and E = 10V, I = 200 A, η= 0.9, λ= 0.028, θp = 730 °C, θ0 = 25 °C, v = 5 mm/s, κ= λ/ρcp = 0.028/0.044:

730 − 25 =

 −5( wp + rp )  0.9 x 20 x 200 / 5 x exp   2π x0.028rp  2 x0.028 / 0.0044 

-1.76 + LNRP = -0.393 (WP + RP)

(EQ 58) (EQ 59)

Solving Eq 21 using the above values, by iteration: •

•

•

ASSUME RP = 3.4 AND PUT THIS VALUE OF RP INTO EQ 59 TO OBTAIN WP = -2.036. SUBSTITUTE RP = 3.4 AND WP = -2.036 INTO EQ 55 AND SOLVE FOR YP: YP = 2.7. SUBSTITUTE YP AND WP INTO EQ 56: THE LEFT SIDE OF EQ 56 = 0.176; THE RIGHT SIDE OF EQ 56 = 0.158. THE RESULTS DO NOT SATISFY EQ 56. ASSUME RP = 3.5 AND PUT THIS VALUE OF RP INTO EQ 59 TO OBTAIN WP = -2.2. SUBSTITUTE RP = 3.5 AND WP = -2.2 INTO EQ 55 AND SOLVE FOR YP: YP = 2.72. SUBSTITUTE YP AND WP INTO EQ 56: THE LEFT SIDE OF EQ 56 = 0.179; THE RIGHT SIDE OF EQ 56 = 0.146. THE RESULTS DO NOT SATISFY EQ 56. ASSUME RP = 3.3 AND PUT THIS VALUE OF RP INTO EQ 59 TO OBTAIN WP = -1.86. SUBSTITUTE RP = 3.3 AND WP = -1.86 INTO EQ 55 AND SOLVE FOR YP: YP = 2.726. SUBSTITUTE YP AND WP INTO EQ 56: THE LEFT SIDE OF EQ 56 = 0.170; THE RIGHT SIDE OF EQ 56 = 0.171. EQUATION 56 IS NOW SATISFIED.

The full HAZ width equals 2yp = 2 × 2.726 = 5.5 mm. Comparing this result, 5.5 mm, with the result obtained with Eq 48 (that is, 5.9 mm), we know that Eq 48 can be used to obtain accurate results when calculating HAZ width and peak temperature.

Effect of Tempering Temperature on Quenched and Tempered (Q&T) Steels. If the plate had been quenched

and then tempered to 430 °C (810 °F), then any region heated above that temperature will have been "over-tempered" and may exhibit modified properties. It would then be reasonable to consider the modified zone as being "heat affected," with its outer extremity located where Tp = 430 °C (810 °F):

4.13(0.0044)5Yz 1 1 = + 430 − 25 720 1510 − 25

(EQ 60)

resulting in a value for Yz of 14.2 mm (0.568 in.). Effect of Preheating Temperature on Q&T Steels. Assume that the Q&T steel described above was preheated to a

temperature, T0, of 200 °C (390 °F):

4.13(0.0044)5Yz 1 1 = + 430 − 200 720 1510 − 200

(EQ 61)

resulting in a value for Yz of 28.4 mm > 14.2 mm (1.14 > 0.568 in.). Therefore, increasing the preheating temperature will increase the value of Yz. Effect of Energy Input on Q&T Steels. Assume that the energy input into the Q&T steel (without preheating)

increases 50% (that is, 1.08 kJ/mm, or 27.4 kJ/in.):

4.13(0.0044)5Yz 1 1 = + 430 − 200 1080 1510 − 25

(EQ 62)

resulting in a value for Yz of 21.3 mm > 14.2 mm (0.839 > 0.568 in.). Therefore, increasing the energy input will increase the value of Yz. Example 2: Welding of 6 mm (0.24 in.) Thick Low-Carbon Steels. The thermal properties needed for heat-flow analysis are assumed to be:

MELTING TEMPERATURE (TM), °C (°F) AUSTENIZATION TEMPERATURE, °C (°F) THERMAL CONDUCTIVITY ( λ ), W/M · K (J/MM · S · °C) VOLUMETRIC SPECIFIC HEAT ( ρ CP), J/MM3 · °C HEAT OF FUSION (L), J/MM3 The welding condition is assumed to be:

CURRENT (I), A ARC VOLTAGE (E), V PREHEAT (T0), °C (°F) HEAT-TRANSFER EFFICIENCY (η)

300 25 25 (77) 0.9

1510 (2750) 730 (1346) 11.7 (0.028) 0.0044 2

Critical Cooling Rate at 550 °C (1020 °F) (Tc). A critical cooling rate exists for each steel composition. If the actual

cooling rate in the weld metal exceeds this critical value, then hard martensitic structures may develop in the HAZ, and there is a great risk of cracking under the influence of thermal stresses in the presence of hydrogen. The best way to determine the critical cooling rate is to make a series of bead-on-plate weld passes in which all parameters, except the arc travel speed, are held constant. After the hardness tests on the weld passes deposited at travel speeds of 6, 7, 8, 9 and 10 mm/s (0.23, 0.28, 0.32, 0.35, and 0.39 in./s), it was found that at the latter two travel speeds, the weld HAZ had the highest hardness. Therefore, the critical cooling rate was encountered at a travel speed of approximately 8 mm/s (0.32 in.s). At this speed, the net energy input is:

H net =

25(300)0.9 = 8.43.75 J / mm 8

(EQ 63)

From Eq 51, the relative plate thickness is:

τ =6

0.0044(550 − 25) = 0.31 843.75

(EQ 64)

Because τ is less than 0.75, the thin-plate equation (Eq 50) applies: 2

R 3  6  = 0.0044   ( 550 − 25 ) = 32.2 2πλ  843.75 

(EQ 65)

resulting in R being equal to 2π(0.028)32.2, which is equal to 5.7 °C/s (10.3 °F/s). This value is the maximum safe cooling rate for this steel and the actual cooling rate cannot exceed this value. Preheating Temperature Requirement. Although the critical cooling rate cannot be exceeded, in the actual welding

operation a preheat can be used to reduce the cooling rate to 5.7 °C/s (10.3 °F/s). Assume that the welding condition is:

CURRENT (I), A ARC VOLTAGE (E), V HEAT-TRANSFER EFFICIENCY (η) TRAVEL VELOCITY (V), MM/S (IN./S) PLATE THICKNESS (T), MM (IN.)

250 25 0.9 7 (0.3) 9 (0.4)

The energy heat input, Hnet, is:

H net =

25(250)0.9 = 804 J / mm 7

(EQ 66)

Assuming that the thin-plate equation (Eq 50) applies: 2

3  R   9    = 32.2 = 0.0044   ( 550 − T0 )  2πλ  max  804 

(EQ 67)

resulting in a T0 of 162 °C (325 °F). The relative plate thickness should be checked:

τ =9

0.0044(550 − 162) = 0.41 804

(EQ 68)

Because τ is less than 0.75, the thin-plate equation does apply. If the initial plate temperature is raised either to or above 162 °C (325 °F), then the cooling rate will not exceed 5.7 °C/s (10.3 °F/s). Effect of Joint Thickness. If the plate thickness increases from 9 to 25 mm (0.36 to 1 in.), but there is the same level of

energy input, then the calculation of the initial plate temperature would be as follows. First, using the thin-plate equation (Eq 50): 2

3  R   25    = 32.2 = 0.0044   ( 550 − T0 )  2πλ  max  804 

(EQ 69)

resulting in a value for T0 of 354 °C (670 °F). The relative plate thickness, τ , should be checked:

τ = 25

0.0044(550 − 354) = 0.82 804

(EQ 70)

Because τ is greater than 0.75, the use of the thin-plate equation is inadequate. Using the thick-plate equation (Eq 49):

( 550 − T0 ) 32.2 =

2

(EQ 71)

804

resulting in a value for T0 of 389 °C (730 °F). The relative plate thickness should be checked:

τ = 25

0.0044(550 − 389) = 0.74 804

(EQ 72)

Although τ is less than, but near to, 0.75, using the thin-plate equation is adequate. Therefore, the initial temperature should be raised to 389 °C (730 °F) to avoid exceeding the cooling rate of 5.7 °C/s (10.3 °F). Now, if the plate thickness increases to 50 mm (2 in.), but there is the same level of energy input, then the thick-plate equation (Eq 49) applies and, again, the value for T0 is 389 °C (730 °F). The relative plate thickness should be checked:

τ = 50

0.0044(550 − 389) = 1, 48 804

Because τ is greater than 0.75, the use of the thick-plate equation is adequate.

(EQ 73)

Under some welding conditions, it is not necessary to reduce the cooling rate by using a preheat. For example, if the plate thickness is 5 mm (0.2 in.) and there is the same level of energy input: 2

3  R   5    = 32.2 = 0.0044   ( 550 − T0 )  2πλ  max  804 

(EQ 74)

resulting in a value for T0 of -24 °C (-11 °F). Therefore, using a preheat is unnecessary. Fillet-Welded "T" Joints. For a weld with a higher number of paths, as occurs in fillet-welded "T" joints, it is sometime

necessary to modify the cooling-rate equation, because the cooling of a weld depends on the available paths for conducting heat into the surrounding cold base metal. When joining 9 mm (0.35 in.) thick plate, where Hnet = 804 J/mm (20.4 kJ/in.), and when there are three legs instead of two, the cooling-rate equation is modified by reducing the effective energy input by a factor of

HNET =

2 3

(804) = 536 J/MM

2 : 3

(EQ 75)

Using the thin-plate equation (Eq 50): 2

3  R   9    = 32.2 = 0.0044   ( 550 − T0 )  2πλ  max  536 

(EQ 76)

resulting in a value for T0 of 254 °C (490 °F). The relative plate thickness should be checked:

τ =9

0.0044(550 − 254) = 0, 44 536

(EQ 77)

Because τ is less than 0.75, using the thin-plate equation is adequate. Therefore, a higher preheat temperature is more necessary than a butt weld because of the enhanced cooling. Example 3: Cooling Rate for the Location at Distance y (in cm) from the Centerline. For a steel plate of 25 mm (1 in.) thickness, (t), the welding condition is assumed to be:

HEAT INPUT (ηEI), KW (CAL/S) TRAVEL SPEED (V), CM/S (IN./S) PREHEAT (T0), °C (°F) NET ENERGY INPUT (HNET), CAL/CM

7.5 (1800) 0.1 (0.04) 20 (68) 18,000

The thermal properties needed for heat flow analysis are assumed to be:

MELTING TEMPERATURE (TM), °C (°F) THERMAL CONDUCTIVITY (λ), W/M · K (CALIT/CM · S · °C) SPECIFIC HEAT (CP), J/KG · C (BTU/LB · °F) DENSITY (ρ), G/CM3 (LB/IN.3)

1400 (2550) 43.1 (0.103) 473 (0.113) 7.8 (0.28)

Assume that one is interested in the critical cooling rate at the location on the surface (z = 0) at distance y = 2 cm from the centerline at the instant when the metal passes through the specific temperature of 615 °C (1140 °F). Initially, the relative plate thickness should be checked. From Eq 51, the relative plate thickness (using English units) is:

τ = 25

7.8 x0.113(615 − 20 = 4, 267 18, 000

Because τ is greater than 0.75, this plate can be treated as a thick plate. From Eq 78, cooling rate for thick plate at the location where the variables are w and r and at critical temperature θ= θc is:

∂θ ∂θ −vη EI  −v ( w + r )   − w v  w   = −v = exp   x  r ² − 2κ 1 + r   ∂t ∂w 2πλ r  2κ  

(EQ 78)

To solve Eq 78, we need to calculate the value of w and r first. From Eq 21, the temperature distribution of thick plate, and Eq 36 where z = 0, we can get:

θ − θ0 =

η EI  −v ( w + r )  exp   2πλ r 2κ  

(EQ 79)

and r= w²+y² By substituting the welding condition and material properties into Eq 21, Eq 36, and r = following simultaneous equations:

615 − 20 =

w² + y ² , we can obtain the

1800  −0.1( w + r )  exp  22π x0.103r  2 x0.117 

and r= w²+4 The value of w and r can be solved by using iteration techniques to solve the above simultaneous equation. The result is that w = -3 cm and r = 3.606 cm. Substituting w and r into Eq 78:

∂θ ∂θ −0.1x1800 0.1  −3    −0.1(−3 + 3.606   −(−3) = −v = x exp  x − 1 +   ∂t θ C ∂w 2π x0.103x3.606 2 x0.117    3.606² 2 x0.117  3.606   Therefore:

∂θ ∂t

= −9.78°C / s θC

From Eq 49 we can calculate the cooling rate along the centerline at the same temperature (615 °C, or 1140 °F):

R=

2π x0.103(615 − 20)² = 12.7°C / s 18, 000

Also from Eq 32, we can calculate the cooling rate in the heat-affected zone at a temperature of 615 °C (1140 °F):

  615 − 20 1.7  2  ∂θ −1  0.9843 − 0.486    x 1 tan = 0.35   +       0.125 ∂t θ C =615      π   336 / 2.362 

0.8

= 4°C / s

Therefore, at the same temperature, the cooling rate at the centerline is greater than the cooling rate at the location a distance y from the centerline. In addition, the cooling rate of the heat-affected zone is less than the cooling rate in the weld pool at the same temperature. Example 4: Solidification Rate. A weld pass of 800 J/mm (20.3 kJ/in.) in net energy input is deposited on a steel plate. The initial temperature is 25 °C (75 °F). The solidification time would be:

St =

2(800) = 0.94( s ) 2π (0.028)(0.0044)(1510 − 25)²

(EQ 80)


22. HEAT FLOW IN WELDING, CHAPTER 3, WELDING HANDBOOK, VOL 1, 7TH ED., AWS, 1976 23. C.M. ADAMS, JR. COOLING RATE AND PEAK TEMPERATURE IN FUSION WELDING, WELD. J., VOL 37 (NO. 5), 1958, P 210S-215S 24. C.M. ADAMS, JR., COOLING RATES AND PEAK TEMPERATURES IN FUSION WELDING, WELD. J., VOL 37 (NO. 5), P 210-S TO 215-S 25. H. KIHARA, H. SUZUKI, AND H. TAMURA, RESEARCH ON WELDABLE HIGH-STRENGTH STEELS, 60TH ANNIVERSARY SERIES, VOL 1, SOCIETY OF NAVAL ARCHITECTS OF JAPAN, TOKYO, 1957 26. C.E. JACKSON, DEPARTMENT OF WELDING ENGINEERING, THE OHIO STATE UNIVERSITY LECTURE NOTE, 1977 27. C.E. JACKSON AND W.J. GOODWIN, EFFECTS OF VARIATIONS IN WELDING TECHNIQUE ON THE TRANSITION BEHAVIOR OF WELDED SPECIMENS--PART II, WELD. J., MAY 1948, P 253-S TO 266-S


Parametric Effects

To show the effects of material property and welding condition on the temperature distribution of weldments, the welding of 304 stainless steel, low-carbon steel, and aluminum are simulated for three different welding speeds: 1.0, 5.0, and 8.0 mm/s (0.04, 0.02, and 0.03 in./s). The thermal conductivity and thermal diffusivity of 304 stainless steel are 26 W/m · K (0.062 cal/cm · s · °C) and 4.6 mm2/s (0.007 in.2/s), respectively. For low-carbon steel, the respective values are 50 W/m K (0.12 cal/cm s °C) and 7.5 mm2/s (0.012 in.2/s), whereas for aluminum, the respective values are 347 W/m K (0.93 cal/cm · s · °C) and 80 mm2 (0.12 in.2/s). The heat input per unit weld length was kept constant, 4.2 kJ/s (1 kcal/s), for all cases. The parametric results are described below. Effect of Material Type. Figures 5(a), 5(b), and 5(c) depict the effect of thermal properties on isotemperature contours

for a heat input of 4.2 kJ/s (1 kcal/s) and travel speeds of 1.0, 5.0, and 8.0 mm/s (0.04, 0.02, and 0.3 in./s). The temperature spreads over a larger area and causes a larger weld pool (larger weld bead) for low-conductivity material. The isotemperature contours also elongate more toward the back of the arc for low-conductivity material. For aluminum, a larger heat input would be required to obtain the same weld size as the stainless steel weldment.

FIG. 5(A) EFFECT OF THERMAL PROPERTY ON ISOTEMPERATURE CONTOURS FOR A HEAT INPUT OF 4.2 KJ/S (1000 CAL/S) AT A WELDING SPEED, V, OF 1 MM/S (0.04 IN./S) AND THE RESPECTIVE THERMAL CONDUCTIVITIES OF EACH MATERIAL (REFER TO TEXT FOR VALUES). VALUES FOR X AND Y ARE GIVEN IN CM, AND TEMPERATURES ARE GIVEN IN °C.

FIG. 5(B) EFFECT OF THERMAL PROPERTY ON ISOTEMPERATURE CONTOURS FOR A HEAT INPUT OF 4.2 KJ/S (1000 CAL/S) AT A WELDING SPEED, V, OF 5 MM/S (0.02 IN./S) AND THE RESPECTIVE THERMAL CONDUCTIVITIES OF EACH MATERIAL (REFER TO TEXT FOR VALUES). VALUES FOR X AND Y ARE GIVEN IN CM, AND TEMPERATURES ARE GIVEN IN °C.

FIG. 5(C) EFFECT OF THERMAL PROPERTY ON ISOTEMPERATURE CONTOURS FOR A HEAT INPUT OF 4.2 KJ/S (1000 CAL/S) AT A WELDING SPEED, V, OF 8 MM/S (0.3 IN./S) AND THE RESPECTIVE THERMAL CONDUCTIVITIES OF EACH MATERIAL (REFER TO TEXT FOR VALUES). VALUES FOR X AND Y ARE GIVEN IN CM, AND TEMPERATURES ARE GIVEN IN °C.

Welding Speed. Figures 6(a), 6(b), and 6(c) show the effect of welding speed on isotemperature contours. When the

travel speed increases, the weld size decreases and the isotemperature contours are more elongated toward the back of the arc. Larger heat inputs would be required for faster travel speeds in order to obtain the same weld size.

FIG. 6(A) EFFECT OF WELDING SPEED, V, ON ISOTEMPERATURE CONTOURS OF 304 STAINLESS STEEL FOR 4.2 KJ/S (1000 CAL/S) HEAT INPUT

FIG. 6(B) EFFECT OF WELDING SPEED, V, ON ISOTEMPERATURE CONTOURS OF LOW-CARBON STEEL FOR 4.2 KJ/S (1000 CAL/S) HEAT INPUT

FIG. 6(C) EFFECT OF WELDING SPEED, V, ON ISOTEMPERATURE CONTOURS OF ALUMINUM FOR 4.2 KJ/S (1000 CAL/S) HEAT INPUT

Figure 7 shows the effect of welding heat input on the peak temperature at two locations, 6.4 mm (

1 1 in.) and 13 mm ( 4 2

in.) from the weld centerline. (The material simulated in this illustration is low-carbon steel plate with a large thickness.) Within the practical range of welding conditions for the GMAW process of low-carbon steels, the peak temperature at both locations increases linearly, as the welding current increases, and decreases exponentially, with the travel speed. To •

increase welding current proportionally with travel speed for a constant heat input per unit weld length (that is. Q /V = constant), the peak temperature at two locations increases with the travel speed, which implies a larger weld size resulting from higher heat input. The welding current has a more significant effect on peak temperature than does the travel speed. However, the increase in peak temperature is more dominant at the 6.4 mm (

1 in.) location. The influence of 4

proportional increase in welding current and travel speed on peak temperature diminishes as the distance from the weld centerline increase and when the travel speed becomes high.

FIG. 7 PEAK TEMPERATURE DETERMINED BY THE POINT HEAT SOURCE SOLUTION. (A) AT 6.4 MM ( FROM THE CENTER. (B) AT 13 MM (

1 IN.) 2

1 IN.) FROM THE CENTER 2


Thermophysical Properties of Selected Engineering Materials For the sake of convenience when using heat-flow equations, the thermal properties of selected engineering materials are provided in Table 2.

TABLE 2 THERMAL PROPERTIES OF SELECTED ENGINEERING ALLOYS ALLOY

CARBON STEEL ~0.5% C

DENSITY (ρ), AT 20 °C (68 °F), G/CM3

SPECIFIC HEAT (CP), AT 20 °C (68 °F), KJ/KG ·K

THERMAL CONDUCTIVITY (κ) AT 20 °C (68 °F, W/M · K)

COEFFICIENT OF THERMAL EXPANSION (α), AT 20 °C (68 °F), 10 -5/°C

7.833

0.465

54

1.474

~1.0% C ~1.5% C ALUMINUM AL-CU (94-96 % AL, 3-5% CU) AL-SI (86.5% AL. 1% CU) AL-SI (78-80% AL, 20-22% SI) AL-MG-SI (97% AL, 1% MG, 1% SI, 1% MN) TITANIUM ALLOY

STAINLESS STEEL CHROMIUM-NICKEL AUSTENITIC CHROMIUM FERRITIC CHROMIUM MARTENSITIC CARBON STEEL

7.801 7.753

0.473 0.486

43 36

1.172 0.970

2.787 2.659 2.627 2.707

0.883 0.867 0.854 0.892

164 137 161 177

6.676 5.933 7.172 7.311

4.500

0.52

16

0.84

DENSITY (ρ), AT 100 °C (212 °F), G/CM2

SPECIFIC HEAT (CP), AT 0-100 °C (32-212 °F), KJ/KG · K

THERMAL CONDUCTIVITY (κ), AT 100 °C (212 °F), W/M · K

COEFFICIENT OF THERMAL EXPANSION (α), AT 0538 °C (32-1000 °F), 10 -5/°C

7.800-8.000

0.46-0.50

18.7-22.8

1.700-1.920

7.800 7.800

0.46-0.50 0.42-0.46

24.4-26.3 28.7

1.120-1.210 1.160-1.210

7.800

0.48

60

1.17


References

1. N.S. BOULTON AND H.E. LANCE-MARTIN, RESIDUAL STRESSES IN ARC WELDED PLATES, PROC. INST. MECH. ENG., VOL 33, 1986, P 295 2. D. ROSENTHAL AND R. SCHMERBER, THERMAL STUDY OF ARC WELDING, WELD. J., VOL 17 (NO. 4), 1983, P 2S 3. N.N. RYKALIN, "CALCULATIONS OF THERMAL PROCESSES IN WELDING," MASHGIZ, MOSCOW, 1951 4. K. MASUBUCHI, ANALYSIS OF WELDED STRUCTURES, PERGAMON PRESS, 1980 5. C.L. TSAI, "PARAMETRIC STUDY ON COOLING PHENOMENA IN UNDERWATER WELDING," PH.D THESIS, MIT, 1977 6. C.L. TSAI, FINITE SOURCE THEORY, MODELING OF CASTING AND WELDING PROCESSES II, ENGINEERING FOUNDATION MEETING, NEW ENGLAND COLLEGE (HENNIKER, NH), 31 JULY TO 5 AUG 1983, P 329 7. R. PAVELIC, R. TANAKUCHI, O. CZEHARA, AND P. MYERS, EXPERIMENTAL AND COMPUTED TEMPERATURE HISTORIES IN GAS TUNGSTEN ARC WELDING IN THIN PLATES, WELD. J., VOL 48 (NO. 7), 1969, P 295S 8. S. KOU, 3-DIMENSIONAL HEAT FLOW DURING FUSION WELDING, PROC. OF METALLURGICAL SOCIETY OF AIME, AUG 1980, P 129-138 9. P.G. KOGAN, THE TEMPERATURE FIELD IN THE WELD ZONE, AVE. SVARKA, VOL 4 (NO. 9), 1979, P 8 10. G.M. ECER, H.D. DOWNS, H.D. BRODY, AND M.A. GOKHALE, HEAT FLOW SIMULATION OF PULSED CURRENT GAS TUNGSTEN ARC WELDING, MODELING OF CASTING AND WELDING PROCESSES, ENGINEERING FOUNDATION 1980 MEETING (RINDGE, NH), 3-8 AUG 1980, P 139160

11. H. HIBBITT AND P. MARCAL, A NUMERICAL THERMOMECHANICAL MODEL FOR WELDING AND SUBSEQUENT LOADING OF A FABRICATED STRUCTURE, COMPUT. STRUCT., VOL 3, 1973, P 1145 12. E. FRIEDMAN, THERMOMECHANICAL ANALYSIS OF THE WELDING PROCESS USING FINITE ELEMENT METHODS, TRANS. ASME, AUG 1975, P 206 13. Z. PALEY AND P. HIBBERT, COMPUTATION OF TEMPERATURE IN ACTUAL WELD DESIGN, WELD. J., VOL 54 (NO. 11), 1975, P 385.S 14. T. NAKA, TEMPERATURE DISTRIBUTION DURING WELDING, J. JPN. WELD. SOC., VOL 11 (NO. 1), 1941, P 4 15. K. MASUBUCHI AND T. KUSUDA, TEMPERATURE DISTRIBUTION OF WELDED PLATES, J. JPN WELD. SOC., VOL 22 (NO. 5), 1953, P 14 16. C.L. TSAI AND C.A. HOU, THEORETICAL ANALYSIS OF WELD POLL BEHAVIOR IN THE PULSED CURRENT GTAW PROCESS, TRANSPORT PHENOMENA IN MATERIALS PROCESSING, ASME WINTER ANNUAL MEETING, 1983 17. E. FRIEDMAN, "FINITE ELEMENT ANALYSIS OF ARC WELDING," REPORT WAPD-TM-1438, DEPARTMENT OF ENERGY, 1980 18. J.S. FAN AND C.L. TSAI, "FINITE ELEMENT ANALYSIS OF WELDING THERMAL BEHAVIOR IN TRANSIENT CONDITIONS," 84-HT-80, ASME 19. R.L. APPS AND D.R. MILNER, HEAT FLOW IN ARGON-ARC WELDING, BR. WELD. J., VOL 2 (NO. 10), 1955, P 475 20. H.S. CARSLAW AND J.C. JAEGER, CONDUCTION OF HEAT IN SOLIDS, OXFORD PRESS 21. C.L. TSAI AND J.H. WU, "AN INVESTIGATION OF HEAT TRANSPORT PHENOMENA IN UNDERWATER WELDING," PRESENTED AT THE ASME WINTER ANNUAL MEETING (MIAMI BEACH, FL), 1985 22. HEAT FLOW IN WELDING, CHAPTER 3, WELDING HANDBOOK, VOL 1, 7TH ED., AWS, 1976 23. C.M. ADAMS, JR. COOLING RATE AND PEAK TEMPERATURE IN FUSION WELDING, WELD. J., VOL 37 (NO. 5), 1958, P 210S-215S 24. C.M. ADAMS, JR., COOLING RATES AND PEAK TEMPERATURES IN FUSION WELDING, WELD. J., VOL 37 (NO. 5), P 210-S TO 215-S 25. H. KIHARA, H. SUZUKI, AND H. TAMURA, RESEARCH ON WELDABLE HIGH-STRENGTH STEELS, 60TH ANNIVERSARY SERIES, VOL 1, SOCIETY OF NAVAL ARCHITECTS OF JAPAN, TOKYO, 1957 26. C.E. JACKSON, DEPARTMENT OF WELDING ENGINEERING, THE OHIO STATE UNIVERSITY LECTURE NOTE, 1977 27. C.E. JACKSON AND W.J. GOODWIN, EFFECTS OF VARIATIONS IN WELDING TECHNIQUE ON THE TRANSITION BEHAVIOR OF WELDED SPECIMENS--PART II, WELD. J., MAY 1948, P 253-S TO 266-S Fluid Flow Phenomena During Welding C.R. Heiple and P. Burgardt, EG&G Rocky Flats

Introduction MOLTEN WELD POOLS are dynamic. Liquid in the weld pool in acted on by several strong forces, which can result in high-velocity fluid motion. Fluid flow velocities exceeding 1 m/s (3.3 ft/s) have been observed in gas tungsten arc (GTA) welds under ordinary welding conditions, and higher velocities have been measured in submerged arc welds. Fluid flow is important because it affects weld shape and is related to the formation of a variety of weld defects. Moving liquid transports heat and often dominates heat transport in the weld pool. Because heat transport by mass flow depends on the

direction and speed of fluid motion, weld pool shape can differ dramatically from that predicted by conductive heat flow. Temperature gradients are also altered by fluid flow, which can affect weld microstructure. A number of defects in GTA welds have been attributed to fluid flow or changes in fluid flow, including lack of penetration, top bead roughness, humped beads, finger penetration, and undercutting. Instabilities in the liquid film around the keyhole in electron beam and laser welds are responsible for the uneven penetration (spiking) characteristic of these types of welds. Fluid Flow Phenomena During Welding C.R. Heiple and P. Burgardt, EG&G Rocky Flats

Mass Transport in the Arc High-velocity gas motion occurs in and around the arc during welding. The gas motion is partially due to cover gas flow, but, more importantly, it is driven by electromagnetic forces associated with the arc itself. In gas metal arc (GMA) welds, liquid filler metal is also being transferred through the arc from the electrode to the workpiece. Both the mode and velocity of metal transfer have major effects on weld pool shape. Mass transport in and around the arc is important in GTA welding (GTAW) and even more so in GMA welding (GMAW); a detailed description of this phenomenon can be found in the articles on arc physics and metal transfer and weld behavior in this Volume. Only the effects of this transport on the weld pool will be discussed in this article. Fluid Flow Phenomena During Welding C.R. Heiple and P. Burgardt, EG&G Rocky Flats

Gas Tungsten Arc Welding Most experimental and theoretical work on weld pool fluid flow and its effects has been directed toward GTAW. The motivation for much of this work was the observation of dramatically different weld pool shapes for GTA welds made using identical welding parameters on different heats of the same material with the same nominal composition. An extreme example of weld shape variability is shown in Fig. 1. Early observations of variable weld shape (often referred to as variable penetration) were not only an intellectual puzzle but also an indication of a growing practical problem. Gas tungsten arc welding is commonly used for high-precision, high-quality automated welding applications, where reproducibility of weld shape or penetration is critical.

FIG. 1 PARTIAL-PENETRATION GTA WELDS MADE UNDER THE SAME WELDING CONDITIONS ON TWO HEATS OF TYPE 304L STAINLESS STEEL HAVING THE SAME NOMINAL COMPOSITION. (A) 3 PPM S, D/W = 0.2. (B) 160 PPM S, D/W = 0.40. 9×

The possibility that fluid flow in the weld pool could alter weld shape has been recognized for many years. For example, in 1965 Christensen et al. (Ref 1) proposed that convection is partially responsible for deviations in weld pool shapes from those predicted by conduction solutions. The forces driving fluid flow in GTA weld pools have also been long known. The four primary driving forces are surface tension gradients, electromagnetic or Lorentz forces, buoyancy forces, and aerodynamic drag forces caused by passage of the arc plasma over the weld pool surface.

Surface-Tension-Driven Fluid Flow Model. Surface tension gradients were first proposed by Ishizake et al. (Ref 2) as

potential driving forces for weld pool fluid flow. Surface-tension-driven fluid flow was first described by Thomson (Ref 3) in 1855, but the phenomenon is commonly called Marangoni convection from the work of Carlo Marangoni (Ref 4). In 1982, Heiple and Roper (Ref 5) proposed that surface tension gradients are commonly the dominant forces driving fluid flow in GTA welds and that these gradients could be drastically altered by very small concentrations of certain trace elements. Surface tension gradients exist on a weld pool surface because the surface tension is temperature dependent, and there are large temperature gradients on a weld pool surface. For pure metals and many alloys, the surface tension decreases as temperature increases; that is, the surface tension temperature coefficient is negative. For weld pools in such materials, the surface tension will be greatest on the coolest part of the pool surface at the edge and lowest on the hottest part under the arc near the center of the pool. Such a surface tension gradient produces outward surface fluid flow, as shown schematically in Fig. 2(a). This fluid flow pattern transfers heat efficiently from the hottest part of the weld pool (near the center) to the edge and produces a relatively wide and shallow weld.

FIG. 2 SCHEMATIC SHOWING SURFACE FLUID FLOW (TOP) AND SUBSURFACE FLUID FLOW (BOTTOM) IN THE WELD POOL. (A) NEGATIVE SURFACE TENSION TEMPERATURE COEFFICIENT (PURE MATERIAL). (B) POSITIVE SURFACE TENSION TEMPERATURE COEFFICIENT (SURFACE-ACTIVE ELEMENTS PRESENT)

Certain elements are surface active in molten metals; that is, they segregate to the surface of the solvent liquid metal and lower the magnitude of the surface tension, often drastically. Small concentrations of surface-active additions can also change the temperature dependence of the surface tension of the solvent metal or alloy so that, for a limited range of temperature above the melting point, the surface tension increases with increasing temperature. With a positive surface tension temperature coefficient, the surface tension will be highest near the center of the weld pool. Such a surface tension gradient will produce fluid flow inward along the surface of the weld pool and then down, as indicated schematically in Fig. 2(b). This fluid flow pattern transfers heat efficiently to the bottom of the weld pool and produces a relatively deep, narrow weld. Experimental Observations. This physical model was developed and verified with a series of experiments in which stainless steel base metal was doped with low concentrations of various elements and the effect of the doping on weld pool shape measured. High-speed motion pictures of the weld pool surface suggested the fluid flow patterns indicated in Fig. 2. The addition of sulfur, oxygen, selenium, and tellurium to stainless steel in low concentrations (less than 150 ppm) was shown to substantially increase GTA weld depth-to-width ratio (d/w). All these elements are known to be highly surface active in liquid iron. Measurements of the temperature dependence of the surface tension for steels with different GTA weld penetration characteristics produced an impressive correlation between a positive surface tension temperature coefficient arising from surface-active impurities and high d/w ratio welds (Fig. 3).

FIG. 3 PLOT OF SURFACE TENSION VERSUS TEMPERATURE FOR TWO LIQUID STEELS. THE DATA LABELED "HIGH D/W HEAT" ARE FROM MATERIAL HAVING APPROXIMATELY 160 PPM MORE SULFUR THAN THE MATERIAL LABELED "LOW D/W HEAT." THE DASHED LINES INDICATE THE EXPECTED BEHAVIOR OF THE SURFACE TENSION ABOVE THE MAXIMUM TEMPERATURE STUDIED. SOURCE: REF 6

When additions were made to the weld pool of elements known to react with surface-active elements already present in the steel to form compounds that are not surface active, the GTA weld d/w ratio decreased. Aluminum reacts with oxygen and produced wider, shallower welds in 21-6-9 stainless steel. Cerium reacts with both sulfur and oxygen and also produced lower d/w ratio welds. The effects of trace elements on weld shape have also been observed in a number of other alloys. These observations are summarized in Table 1.

TABLE 1 EFFECT OF TRACE ELEMENT IMPURITIES ON GTA WELD PENETRATION OF SELECTED ALLOYS

ALLOY SYSTEM

ZIRCALOY-2 IRON-BASE ALLOYS STAINLESS STEELS 304, 316, 21-6-9 JBK-75 AISI 8630 2.5CR1MO NICKEL-BASE ALLOYS INCONEL 600, 718 (A) (B)

TRACE ELEMENT IMPURITY INCREASES DECREASES WELD PENETRATION WELD PENETRATION CHLORINE ... SULFUR, OXYGEN, CALCIUM, ALUMINUM, CERIUM(A), SELENIUM(A), LANTHANUM(A), SILICON(B), TELLURIUM(A) TITANIUM(B)

SULFUR

OXYGEN

UNCOMMON IMPURITIES EFFECT NEGLIGIBLE OR UNCERTAIN

The simple physical model illustrated in Fig. 2 has been remarkably successful in qualitatively explaining trace element effects. For example, changing GTA welding conditions alters the magnitude and distribution of arc energy input to the weld, which in turn changes temperature gradients on the weld pool surface. From Fig. 2, a change in welding conditions that makes the center of the weld hotter, such as increasing current, should drive the existing fluid flow pattern more strongly. As shown in Fig. 4, increasing current improves the d/w ratio of steel doped with surface-active elements and reduces it for high-purity base metal. If the center of the weld becomes so hot that there is a region where the temperature coefficient of the surface tension is no longer positive, then the fluid flow pattern necessary for deep penetration is disrupted and the d/w ratio decreases. This effect is seen at high currents in Fig. 4. Similar results have been obtained for other welding parameters. The surface temperature at which the change from positive to negative surface tension

temperature coefficient occurs for stainless steel is estimated by extrapolation in Fig. 3 to be about 2050 °C (3720 °F). Detailed thermodynamic calculations of the temperature dependence of the surface tension of iron-sulfur alloys predict that the transition from positive to negative dγ/dT will occur at 2032 °C (3690 °F) for the high-sulfur alloy in Fig. 3. Recent spectrographic weld pool temperature measurements and numerical simulations have indicated that this temperature can be exceeded in stainless steel GTA weld pools under normal welding conditions.

FIG. 4 PLOT OF WELD D/W RATIO VERSUS WELD CURRENT FOR THE STARTING BASE METAL (TYPE 304 STAINLESS STEEL WITH VERY LOW RESIDUAL IMPURITY CONTENT) AS WELL AS FOR ZONES DOPED WITH SULFUR AND SELENIUM. SOURCE: REF 7

The surface-tension-driven fluid flow model should be applicable to non-arc processes, provided the energy input distribution is similar to a GTA arc. This condition is satisfied for conduction-mode electron beam and laser welds. Dramatic increases in weld d/w ratio in selenium-doped zones in stainless steel have been observed for both traveling laser and electron beam conduction-mode welds. The weld shape changes were similar to those observed for GTA welds. Conduction-mode electron beam welds can also be used to demonstrate that variations in weld shape with changes in welding parameters, as illustrated in Fig. 4, are not a result of some complex arc/weld pool interaction. One of the results of an investigation of the effect of changes in beam focus on weld shape in electron beam welds on low- and high-sulfur materials is shown in Fig. 5. The high-sulfur material exhibits a maximum in d/w ratio with increasing power density at a moderate power density away from sharp focus, which is analogous to that shown with increasing current in Fig. 4. (The low- and high-sulfur stainless steels have deep and nearly identical penetration near sharp focus. Near sharp focus, where the power density is high, penetration is by a keyhole mechanism in which average penetration is unrelated to surfacetension-driven fluid flow.) Measurements of the electron beam power-density distribution verified that there were no anomalous changes in the beam, such as a beam width maximum, with increasing peak power density. The d/w ratio maximum away from sharp focus is therefore proposed to originate from exactly the same mechanism as for increasing current with GTA welds.

FIG. 5 PLOT OF ELECTRON BEAM WELD POOL RATIO (D/W) VERSUS ELECTRON BEAM POWER DENSITY FOR LOW-SULFUR (20 PPM) AND HIGH-SULFUR (>120 PPM) TYPE 304L STAINLESS STEEL. KEYHOLE FORMATION BEGINS AT ABOUT 2 × 103 W/MM2.

Measurements of electron beam power distribution were made as a function of beam focus and were used to calculate the beam power density at the d/w maximum. The power density was also calculated at the d/w maximum for GTA welds. Calculation of weld pool surface temperatures using a traveling distributed heat source model showed the central surface temperatures to be essentially identical at the d/w maxima for the electron beam and GTA processes. The calculated surface temperatures using the traveling distributed heat source conduction approximation are much too high because, as indicated previously, most of the heat transport in the weld pool is by mass flow rather than by conduction. Nevertheless, the equality of the calculated weld pool peak temperatures at the d/w maxima provides strong confirmation that the mechanism responsible for the presence of the maximum in weld d/w with increasing input power density is independent of heat source and is not a result of an arc phenomenon. A final example of the success of the fluid flow model in explaining GTA welding phenomena is provided by butt welding together two steels with large differences in weld penetration characteristics. The weld pool is not centered over the joint; rather, it is displaced toward the material with low d/w behavior (Ref 8), as indicated schematically in Fig. 6. The low d/w material has a low concentration of surface-active impurities and therefore a high surface tension. Thus, there is a net surface tension gradient across the weld pool toward the low d/w material, producing the fluid flow pattern and weld cross section indicated in Fig. 6. The actual fluid flow pattern is certain to be more complicated than that illustrated.

FIG. 6 SCHEMATIC SHOWING TYPICAL FLUID FLOW GENERATED WHEN BUTT WELDING TWO HEATS OF MATERIAL WITH DIFFERENT PENETRATION CHARACTERISTICS. SOURCE: REF 8

Numerical Simulations. Direct measurement of weld pool fluid flow is very difficult. Surface flow has been studied by

observing the motion on the surface of slag or intentionally added particles, as has bulk flow in transparent liquid/solid systems using simulated welding heat sources. There are a number of numerical calculations of weld pool fluid flow and shape, the most comprehensive of which appears to be that of Zacharia et al. (Ref 9). These numerical fluid flow simulations agree in broad terms with the physical model illustrated in Fig. 2; however, they differ in detail with Fig. 2 and with one another. One result of the analysis of Zacharia is that the weld pool surface temperature reached in stationary welds exceeds that in traveling welds. Trace element effects should therefore be less pronounced in stationary welds, because a larger portion of the pool surface is above the temperature range where a positive surface tension temperature coefficient exists. This prediction is in agreement with experimental observations. Interactions. Experiments in which low concentrations of various elements have been added to the base metal have

demonstrated clear changes in weld pool shape. However, when weld d/w ratio is plotted versus chemical analysis for numerous heats of material, considerable scatter is observed. A comparison of weld d/w ratio in a standard weldability test versus sulfur content for about 200 lots of steel is shown in Fig. 7. There is a clear trend of increasing d/w with increasing sulfur, but the variability for a given sulfur content is substantial. Some of this variability is associated with imprecision in chemical analysis for sulfur and some with variations in oxygen content. However, it appears that much of the variability is related to interactions of the surface-active elements sulfur and oxygen with other components of the steel. Calcium is known to react with oxygen and, to a lesser extent, with sulfur to form stable compounds unlikely to be surface active. Aluminum and silicon also react with oxygen to form stable compounds. Thus, the amount of sulfur and oxygen available for segregation to the weld pool surface is a complicated function of the total weld pool chemistry.

FIG. 7 PLOT OF WELD D/W RATIO VERSUS SULFUR CONTENT FOR APPROXIMATELY 200 HEATS OF TYPE 304L STAINLESS STEEL. EACH POINT IS AN AVERAGE OF MULTIPLE SULFUR ANALYSES AND WELD D/W RATIO MEASUREMENTS. IF SINGLE VALUES ARE USED, THE SCATTER IS GREATER. SOURCE: REF 10

There is an additional complication, illustrated in Fig. 8. When oxygen is added to the shielding gas, weld d/w ratio increases, passes through a maximum, and then declines with increasing oxygen content. Similar effects are seen with SO2 additions. The surface of welds made with torch gas concentrations above the d/w maximum are heavily oxidized. A possible explanation for the decreasing d/w ratios at higher oxygen concentrations is that a liquid oxide film (slag) is formed on the weld pool surface, altering the surface tension gradients.

FIG. 8 PLOT OF WELD D/W RATIO VERSUS OXYGEN CONCENTRATION IN THE TORCH GAS FOR GTA BEAD-ONPLATE WELDS ON 21-6-9 STAINLESS STEEL. SOURCE: REF 11

Oxygen can be added to the weld pool in other ways. Oxidizing the plate surface prior to welding adds oxygen to the weld pool and increases the weld d/w ratio. Wire brushing and grit blasting an originally clean surface increases both the surface oxide thickness and the surface area, thereby adding oxygen to the weld pool. The d/w ratio of JBK-75 stainless steel (a modification of ASTM A-286) GTA welds as a function of amount of wire brushing is similar in form to Fig. 8. Welds on plate brushed beyond the weld d/w maximum had extensive slag on the weld surface. The weld d/w ratio on JBK-75 stainless steel appears to be particularly sensitive to oxygen additions. Similar wire-brushing experiments on type 304 stainless steel showed smaller effects. Although the origin of the difference in sensitivity between JBK-75 and type 304 is not known, one possibility is indicated by surface tension measurements on iron-silicon alloys in contact with carbon dioxide. The effect of CO2 on the surface tension was a strong function of silicon content. For low-silicon alloys, the surface tension dropped sharply when contacted with CO2. For alloys with more than 1.2% Si, the surface tension increased when contacted with CO2. The different behavior was attributed to differences in slag formation on the liquid metal surface. Thus, silicon appears to interfere with the ability of oxygen to produce a positive surface tension temperature coefficient on liquid iron. The JBK75 stainless steel used in the wire-brushing experiments contained only 0.06% Si, and the 21-6-9 used for the torch gas experiment (Fig. 8) contained only 0.16% Si. Type 304 stainless steel typically contains more than 0.5% Si. Thus, the high sensitivity of JBK-75 to oxygen additions may be related to its unusually low silicon content. Reports in the literature on the effect of oxygen on GTA weld penetration have been somewhat variable; differences in weld pool silicon content may be partially responsible for this variability. Effects of High Current. Electromagnetic (Lorentz) stirring of the weld pool becomes more important at higher

currents. The direction of the Lorentz force produces the deep penetration fluid flow pattern indicated in Fig. 2(b). At sufficiently high current, the Lorentz force dominates other forces that drive fluid flow, and the effects of trace elements become less important. In addition, the plasma jet becomes stronger with increasing current and at high enough currents produces a significant depression of the weld pool surface. The plasma jet forces are resisted by surface tension forces on the weld pool. Halmoy (Ref 12) has shown that for a traveling weld the radial pressure gradient from the plasma jet tends to transport liquid from the front to the rear of the weld (Fig. 9). For a sufficiently strong pressure gradient, pB, the liquid level under the arc may be pushed down to the bottom of the pool, as shown in Fig. 9(c). High-speed motion pictures of the effect of sulfur additions on weld pool fluid flow in 21-6-9 stainless steel showed behavior similar to that in Fig. 9(c) after large sulfur additions. Sulfur additions substantially reduce the weld pool surface tension and thereby increase the effect of the plasma jet.

FIG. 9 EFFECT OF ARC PRESSURE ON THE WELD POOL FOR STATIONARY AND TRAVELING WELDS. (A) V = 0. (B) V > 0, WITH WEAK RADIAL PRESSURE GRADIENT, P1. (C) V > 0, WITH STRONG RADIAL PRESSURE GRADIENT, PB. SOURCE: REF 12

In addition, a surface depression almost certainly changes the energy distribution input to the weld pool by the arc compared with a flat weld pool surface. Finally, under some high-current conditions, a vortex has been observed near the center of weld pools. Strategies for Controlling Poor and Variable Penetration. A number of techniques have been developed to reduce

penetration variability and improve penetration. Sometimes the weld pool shift illustrated in Fig. 6, which occurs when heats with different penetration characteristics are welded together, can be minimized by very tight heat sinking. The sensitivity of weld shape to trace element differences is a function of welding parameters, as illustrated in Fig. 4 for current; thus, combinations of welding parameters can be chosen that minimize heat-to-heat penetration variability. For wire-fed joints, some joint designs are more tolerant of penetration variability than others. For example, joints with thinner, wider lands are less sensitive than narrower grooves with thicker lands. Material control by selection or specification is another approach. It was originally anticipated that stainless steels with less than 20 ppm S would have poor, but consistent, penetration. However, the data illustrated in Fig. 7 demonstrate that penetration variability is substantial even at these very low sulfur levels. Stainless steels with more than about 100 ppm S generally have consistently good penetration. A related approach is to test incoming material for welding characteristics and then select material with desired welding behavior for critical applications. Finally, the weld pool can be doped with enough surface-active elements to ensure good penetration. Sulfur and oxygen are the most practical dopants for ferrous alloys. The addition of oxygen or sulfur dioxide to the torch shielding gas has been shown to improve penetration and reduce penetration variability; however, this approach presents several practical problems. Oxygen can also be added to the weld pool by oxidizing the weld joint or by otherwise increasing the oxygen content of the weld groove surface--by wire brushing, for example. Probably the most useful approach is to dope the weld pool in the wire-fed joints by using a special filler wire. A group of defense contractors had a special heat of type 380L stainless steel produced with a sulfur content in the range of 100 to 150 ppm. The heat was then converted into weld wire. Welds made with this wire on a variety of stainless steel base metals have exhibited consistent and good penetration.


1. N. CHRISTENSEN, V. DE L. DAVIES, AND K. GJERMUNDSEN, DISTRIBUTION OF TEMPERATURES IN ARC WELDING BR. WELD.J., VOL 12, 1965, P 54-75 2. K. ISHIZAKE, K. MURAI, AND Y. KANBE, "PENETRATION IN ARC WELDING AND CONVECTION IN MOLTEN METAL," DOCUMENT 77-66, INTERNATIONAL INSTITUTE OF WELDING, STUDY GROUP 212, 1966 3. J. THOMSON, ON CERTAIN CURIOUS MOTIONS OBSERVABLE AT THE SURFACES OF WINE AND OTHER ALCOHOLIC LIQUORS, PHILOS.MAG., VOL 10, 1855, P 330-333 4. C. MARANGONI, ÜBER DIE AUSBREITUNG DER TROPFEN EINER FLUESSIGKEIT AUF DER OBERFLAESCHE EINER ANDEREN, ANN. PHYS. CHEM., VOL 143 (NO. 7), 1871, P 337-354

5. C.R. HEIPLE AND J.R. ROPER, MECHANISM FOR MINOR ELEMENT EFFECT ON GTA FUSION ZONE GEOMETRY, WELD. J., VOL 61, 1982, P 97S-102S 6. B.J. KEENE, K.C. MILLS, AND R.F. BROOKS, SURFACE PROPERTIES OF LIQUID METALS AND THEIR EFFECTS ON WELDABILITY, MATER. SCI.TECHNOL., VOL 1, 1985, P 568-571 7. P. BURGARDT AND C.R. HEIPLE, INTERACTION BETWEEN IMPURITIES AND WELDING PARAMETERS IN DETERMINING GTA WELD SHAPE, WELD. J., VOL 65, 1986, P 150S-155S 8. M.J. TINKLER, I. GRANT, G. MIZUNO, AND C. GLUCK, WELDING 304L STAINLESS STEEL TUBING HAVING VARIABLE PENETRATION CHARACTERISTICS, THE EFFECTS OF RESIDUAL, IMPURITY, AND MICRO-ALLOYING ELEMENTS ON WELDABILITY AND WELD PROPERTIES, P.H.M. HART, ED., PAPER 29, THE WELDING INSTITUTE, CAMBRIDGE, 1984 9. T. ZACHARIA, S.A. DAVID, J.M. VITEK, AND T. DEBROY, WELD POOL DEVELOPMENT DURING GTA AND LASER BEAM WELDING OF TYPE 304 STAINLESS STEEL, PART 1-THEORETICAL ANALYSIS, WELD. J., VOL 68, 1989, P 499S-509S 10. P. BURGARDT AND R.D. CAMPBELL, CHEMISTRY EFFECTS ON STAINLESS STEEL WELD PENETRATION, FERROUS ALLOY WELDMENTS, TRANS TECH PUBLICATIONS, SWITZERLAND, 1992, P 379-415 11. C.R. HEIPLE, P. BURGARDT, AND J.R. ROPER, THE EFFECT OF TRACE ELEMENTS ON GTA WELD PENETRATION, MODELING OF CASTING AND WELDING PROCESSES II, J.A. DANTZIG AND J.T. BERRY, ED., TMS-AIME, 1984, P 193-205 12. E. HALMOY, THE PRESSURE OF THE ARC ACTING ON THE WELD POOL, ARC PHYSICS AND WELD POOL BEHAVIOR, THE WELDING INSTITUTE, CAMBRIDGE, 1979, P 259-266

Fluid Flow Phenomena During Welding C.R. Heiple and P. Burgardt, EG&G Rocky Flats

Deep-Penetration Electron Beam and Laser Welds Keyhole Formation. A fundamental difference between arc heat sources and electron or laser beam heat sources is that electron and laser beams are capable of delivering heat over a small area at much higher power densities. As the power density of the welding heat source is increased, the peak surface temperature of the weld pool rises. For many metals, vapor pressure rises nearly exponentially with temperature and becomes appreciable (above 4000 Pa, or 30 torr, for an electron beam operating in a vacuum) near 0.8 Tb, where Tb is the material boiling point in degrees Kelvin. As the surface temperature approaches this value, the liquid surface under the power source is depressed by the vapor pressure. As the liquid moves away from the power source, the surface is depressed and a cavity is formed. This is the basic theory of keyhole formation by an electron or laser beam. Fluid Flow in the Keyhole. The keyhole is a cavity having roughly the size and shape of the beam; that is, it is usually

approximately cylindrical. Fluid flow in the thin layer around the cavity plays a major role in determining the behavior of deep-penetration, keyhole-mode electron beam and laser welds. The first flow required for a traveling weld is transport of metal melted at the front wall of the cavity to the rear (where it eventually solidifies). Because of the weld geometry, this liquid must move around the keyhole cavity as a thin, high-velocity layer on the walls of the cavity. The dominant driving force for this motion appears to be surface tension gradients. As in GTAW, these gradients arise because the surface tension is temperature dependent and there is a substantial temperature difference between the front and rear of the cavity. Calculations of the fluid flow have been performed by Wei and Giedt (Ref 13). Their calculations for pure iron predict that liquid metal is moved from the front to the rear of a sharp focus keyhole in a film that is about 0.02 mm (0.0008 in.) thick and is moving at a velocity of about 250 mm/s (10 in./s). The fluid is driven by a temperature difference of about 400 °C (720 °F). These results are only nominal values from the calculations, but they change little with the assumptions used and are representative of electron beam welds in general.

There is an interesting potential problem with this model. As discussed previously, the surface tension temperature coefficient, dγ/dT, can be positive over a limited temperature range in steels and some other alloys if surface-active impurities are present in sufficient quantity. If dγ/dT is positive, then it would appear that molten metal would not be transported around the cavity. The likely resolution of this difficulty is that the liquid dwells slightly longer in the front of the cavity and is heated by the beam above the temperature at which a transition to negative dγ/dT occurs. Some perturbation of the normal fluid flow may occur under these circumstances. Effects of such a perturbation on weld characteristics appear to be uncommon. The only available evidence for such an effect comes from electron beam welds performed on a heat of type 304 stainless steel with very low residual impurity content, except for about 340 ppm S. These welds had very high porosity compared with identical welds in other heats of material. Instability in Keyhole Fluid Flow. In addition to the steady-state flow of liquid around the keyhole, significant

instability in the fluid motion has been seen. For example, Mara et al. (Ref 14) used side-view, self-illuminated X-ray films to show that the beam locations shifts between full penetration and nearly zero penetration in an irregular fashion. High-speed X-ray photographs show that the variable penetration is caused by a lump of metal that sags into the keyhole from high up on the cavity rear wall. The sequence of events that creates weld penetration irregularity was analyzed by Tong and Giedt (Ref 15) and is illustrated schematically in Fig. 10. First, the beam forms a keyhole and produces a lump of very hot displaced metal at the top rear of the traveling cavity (Fig. 10a). After some period of time, the weld reaches the full penetration allowed by heat flow (Fig. 10b). The lump of displaced metal is unstable and eventually falls into the cavity, partially filling it. The beam must now drill through this additional material. The sequence (a) to (c) in Fig. 10 occurs repeatedly, but irregularly, and produces weld penetration variability. When particularly severe, the irregular penetration is called spiking. Spiking also leads to a tendency for voids to become trapped in the root of the weld.

FIG. 10 SCHEMATIC SHOWING KEYHOLE INSTABILITY. (A) KEYHOLE IS FORMED BY HEAT GENERATED BY ELECTRON BEAM. (B) MAXIMUM PENETRATION THAT CAN BE PRODUCED BY HEAT FLOW. (C) LIQUID COOLS, CAUSING IMPENDING COLLAPSE OF DISPLACED METAL. THE KEYHOLE IS FILLED BY A LUMP OF COOLING MATERIAL AT THE END OF (C), WHICH RETURNS THE KEYHOLE TO CONDITION (A) TO RESTART THE SEQUENCE.

A mathematical model that is the basis for understanding this keyhole instability was developed by Giedt et al. (Ref 13, 15). Stability of the keyhole is basically a balance between vapor pressure, which keeps the cavity open, and surface tension, which tries to close the cavity. Vapor pressure and surface tension are both functions of temperature. Vapor pressure increases nearly exponentially with temperature for many materials, while surface tension generally decreases slowly with temperature. Because there is a substantial temperature gradient from the top to the bottom of the keyhole, both vapor pressure and surface tension vary with depth in the keyhole. Direct measurements of cavity wall temperatures for electron beam welds indicate that the temperature at the top is nearly equal to the melting temperature, while the temperature at the bottom is high enough that the vapor pressure there is large. In type 304 stainless steel, the temperature at the bottom of the cavity is about 2200 °C (1200 °F). The data indicate that the temperature profile is insensitive to welding variables and is determined primarily by material properties, that is, by the temperature where the vapor pressure exceeds about 1330 Pa (10 torr).

Based on temperature measurements of a typical keyhole, the vapor pressure decreases approximately exponentially with height from the bottom of the keyhole and approaches zero at the top of the keyhole. The vapor pressure is balanced primarily by surface tension. For pure iron and many other materials, the surface tension decreases approximately linearly with temperature. Thus, the inward surface tension pressure increases slowly from the bottom to the top of the keyhole. The balance of vapor pressure and surface tension is such that the cavity is stable near the bottom, where the vapor pressure tending to expand the cavity exceeds the surface tension tending to collapse it. At the top, the cavity is unstable because the surface tension tending to collapse it exceeds the vapor pressure tending to expand it. Thus, the cavity will always tend to be filled by liquid originating from above some height H in the keyhole, the pressure crossover height. This cavity filling produces weld penetration irregularity. Weld penetration irregularity is a consequence of the nature of the keyhole in deep-penetration welding and will not generally be solved by simple equipment or material modifications. Penetration irregularity is usually less severe if H (as measured from the top) is a small fraction of the keyhole depth. In such cases the volume of the lump is small, and liquid tends to flow into the cavity relatively smoothly. If H is a large fraction of the keyhole depth, then the liquid lump has a relatively large volume and, when it falls into the keyhole at irregular times, it nearly fills the cavity and produces large penetration variations.


13. P.S. WEI AND W.H. GIEDT, SURFACE TENSION GRADIENT-DRIVEN FLOW AROUND AN ELECTRON BEAM WELDING CAVITY, WELD. J.,VOL 64, 1985, P 251S-259S 14. G.L. MARA, E.R. FUNK, R.C. MCMASTER, AND P.E. PENCE, PENETRATION MECHANISMS OF ELECTRON BEAM WELDING AND THE SPIKING PHENOMENON, WELD. J., VOL 53, 1974, P 246S-251S 15. H. TONG AND W.H. GIEDT, A DYNAMIC INTERPRETATION OF ELECTRON BEAM WELDING, WELD. J., VOL 49, 1970, P 259S-266S Fluid Flow Phenomena During Welding C.R. Heiple and P. Burgardt, EG&G Rocky Flats

Gas Metal Arc Welding Fluid flow certainly occurs in GMA weld pools, but reports on the details of its nature and effects are quite limited. In the spray transfer mode, the impact of the stream of droplets from the electrode on the weld pool forms a substantial depression or crater. This is the mechanism responsible for the typical fingerlike penetration observed with argon as the torch gas. Thus, the depth of penetration is primarily dependent on the momentum of the stream of droplets. A plot of penetration versus the momentum of the droplet stream shows an excellent correlation (Fig. 11). Because the major force driving the droplets toward the weld is drag from the plasma jet, the kinematic viscosity (density times viscosity) of the shielding gas is important. For example, argon has a substantially higher kinematic viscosity than helium, so argon drives the droplet stream more effectively, resulting in a deeper, narrower penetration finger than with helium torch gas. The strength of the plasma jet can also be altered by changing the ambient pressure. For GMA welds on aluminum. Amson and Salter (Ref 17) saw a steady decrease in penetration with pressure until at 13,300 Pa (100 torr) there was essentially no penetration at all. Another way to modify the penetration depth is to spread the stream of droplets over the surface of the weld. This can be accomplished by applying a varying transverse magnetic field.

FIG. 11 PLOT OF GMA WELD PENETRATION IN MILD STEEL FOR SPRAY TRANSFER MODE VERSUS MOMENTUM OF DROPLET STREAM. SOURCE: REF 16


16. W.G. ESSERS AND R. WALTER, HEAT TRANSFER AND PENETRATION MECHANISMS WITH GMA AND PLASMA-GMA WELDING, WELD. J., VOL 62, 1981, P 37S-42S 17. J.C. AMSON AND G.R. SALTER, ANALYSIS OF THE GAS-SHIELDED CONSUMABLE METALARC WELDING SYSTEM, BR. WELD. J., VOL 10, 1963, P 472-483 Fluid Flow Phenomena During Welding C.R. Heiple and P. Burgardt, EG&G Rocky Flats

Submerged Arc Welding Currents commonly employed in submerged arc welding (SAW) are much higher than those used in GTAW or GMAW. Submerged arc welding currents often exceed 1000 A. Thus, the electromagnetic or Lorentz force combined with the tendency of the radial pressure gradient in the moving cavity to transport liquid to the rear of the cavity (Fig. 9) are likely the dominant forces driving fluid flow. The generally accepted flow pattern is indicated schematically in Fig. 12. A cavity is formed at the front of the moving weld pool by arc pressure and by the momentum of drops from the rapidly melting electrode. Metal that is melted at the front of the pool flows underneath and on either side of the cavity. At the rear of the pool, the flow reverses and metal flows back toward the cavity along and near the surface. Flow velocities can be very high; Eichhorn and Engel (Ref 19) measured 4 m/s (13 ft/s) at 720 A. The flow pattern was derived from observations of motion of marker elements, added to the weld pool, as determined from subsequent metallographic sections, and from the motion of radioactive tracer additions. The general features illustrated in Fig. 12 appear to be supported by X-ray fluoroscopy observations. Major perturbations in the indicated fluid flow have been demonstrated, including reversal of the flow direction. These perturbations have been associated with weld defects.

FIG. 12 SCHEMATIC SHOWING TYPICAL FLOW PATTERN IN A SUBMERGED ARC WELD POOL. SOURCE: REF 18


18. J.F. LANCASTER, THE PHYSICS OF WELDING, INTERNATIONAL INSTITUTE OF WELDING, PERGAMON PRESS, OXFORD, 1984, P 243 19. F. EICHHORN AND A. ENGEL, "MASS TRANSFER IN THE WELD POOL," DOCUMENT 201-70, INTERNATIONAL INSTITUTE OF WELDING, STUDY GROUP 212, 1970 Fluid Flow Phenomena During Welding C.R. Heiple and P. Burgardt, EG&G Rocky Flats

References

1. N. CHRISTENSEN, V. DE L. DAVIES, AND K. GJERMUNDSEN, DISTRIBUTION OF TEMPERATURES IN ARC WELDING BR. WELD.J., VOL 12, 1965, P 54-75 2. K. ISHIZAKE, K. MURAI, AND Y. KANBE, "PENETRATION IN ARC WELDING AND CONVECTION IN MOLTEN METAL," DOCUMENT 77-66, INTERNATIONAL INSTITUTE OF WELDING, STUDY GROUP 212, 1966 3. J. THOMSON, ON CERTAIN CURIOUS MOTIONS OBSERVABLE AT THE SURFACES OF WINE AND OTHER ALCOHOLIC LIQUORS, PHILOS.MAG., VOL 10, 1855, P 330-333 4. C. MARANGONI, ÜBER DIE AUSBREITUNG DER TROPFEN EINER FLUESSIGKEIT AUF DER OBERFLAESCHE EINER ANDEREN, ANN. PHYS. CHEM., VOL 143 (NO. 7), 1871, P 337-354 5. C.R. HEIPLE AND J.R. ROPER, MECHANISM FOR MINOR ELEMENT EFFECT ON GTA FUSION ZONE GEOMETRY, WELD. J., VOL 61, 1982, P 97S-102S 6. B.J. KEENE, K.C. MILLS, AND R.F. BROOKS, SURFACE PROPERTIES OF LIQUID METALS AND THEIR EFFECTS ON WELDABILITY, MATER. SCI.TECHNOL., VOL 1, 1985, P 568-571 7. P. BURGARDT AND C.R. HEIPLE, INTERACTION BETWEEN IMPURITIES AND WELDING PARAMETERS IN DETERMINING GTA WELD SHAPE, WELD. J., VOL 65, 1986, P 150S-155S 8. M.J. TINKLER, I. GRANT, G. MIZUNO, AND C. GLUCK, WELDING 304L STAINLESS STEEL

TUBING HAVING VARIABLE PENETRATION CHARACTERISTICS, THE EFFECTS OF RESIDUAL, IMPURITY, AND MICRO-ALLOYING ELEMENTS ON WELDABILITY AND WELD PROPERTIES, P.H.M. HART, ED., PAPER 29, THE WELDING INSTITUTE, CAMBRIDGE, 1984 9. T. ZACHARIA, S.A. DAVID, J.M. VITEK, AND T. DEBROY, WELD POOL DEVELOPMENT DURING GTA AND LASER BEAM WELDING OF TYPE 304 STAINLESS STEEL, PART 1-THEORETICAL ANALYSIS, WELD. J., VOL 68, 1989, P 499S-509S 10. P. BURGARDT AND R.D. CAMPBELL, CHEMISTRY EFFECTS ON STAINLESS STEEL WELD PENETRATION, FERROUS ALLOY WELDMENTS, TRANS TECH PUBLICATIONS, SWITZERLAND, 1992, P 379-415 11. C.R. HEIPLE, P. BURGARDT, AND J.R. ROPER, THE EFFECT OF TRACE ELEMENTS ON GTA WELD PENETRATION, MODELING OF CASTING AND WELDING PROCESSES II, J.A. DANTZIG AND J.T. BERRY, ED., TMS-AIME, 1984, P 193-205 12. E. HALMOY, THE PRESSURE OF THE ARC ACTING ON THE WELD POOL, ARC PHYSICS AND WELD POOL BEHAVIOR, THE WELDING INSTITUTE, CAMBRIDGE, 1979, P 259-266 13. P.S. WEI AND W.H. GIEDT, SURFACE TENSION GRADIENT-DRIVEN FLOW AROUND AN ELECTRON BEAM WELDING CAVITY, WELD. J.,VOL 64, 1985, P 251S-259S 14. G.L. MARA, E.R. FUNK, R.C. MCMASTER, AND P.E. PENCE, PENETRATION MECHANISMS OF ELECTRON BEAM WELDING AND THE SPIKING PHENOMENON, WELD. J., VOL 53, 1974, P 246S-251S 15. H. TONG AND W.H. GIEDT, A DYNAMIC INTERPRETATION OF ELECTRON BEAM WELDING, WELD. J., VOL 49, 1970, P 259S-266S 16. W.G. ESSERS AND R. WALTER, HEAT TRANSFER AND PENETRATION MECHANISMS WITH GMA AND PLASMA-GMA WELDING, WELD. J., VOL 62, 1981, P 37S-42S 17. J.C. AMSON AND G.R. SALTER, ANALYSIS OF THE GAS-SHIELDED CONSUMABLE METALARC WELDING SYSTEM, BR. WELD. J., VOL 10, 1963, P 472-483 18. J.F. LANCASTER, THE PHYSICS OF WELDING, INTERNATIONAL INSTITUTE OF WELDING, PERGAMON PRESS, OXFORD, 1984, P 243 19. F. EICHHORN AND A. ENGEL, "MASS TRANSFER IN THE WELD POOL," DOCUMENT 201-70, INTERNATIONAL INSTITUTE OF WELDING, STUDY GROUP 212, 1970 Fluid Flow Phenomena During Welding C.R. Heiple and P. Burgardt, EG&G Rocky Flats

Selected References

• P. BURGARDT, WELDING VARIABLE EFFECTS ON WELD SHAPE IN ELECTRON BEAM WELDING OF STEELS, FERROUS ALLOY WELDMENTS, TRANS TECH PUBLICATIONS, SWITZERLAND, 1992, P 269-328 • P. BURGARDT AND R.D. CAMPBELL, CHEMISTRY EFFECTS ON STAINLESS STEEL WELD PENETRATION, FERROUS ALLOY WELDMENTS, TRANS TECH PUBLICATIONS, SWITZERLAND, 1992, P 379-415 • C.R. HEIPLE AND P. BURGARDT, PENETRATION IN GTA WELDING, WELDABILITY OF MATERIALS, R.A. PATTERSON AND K.W. MAHIN, ED., ASM INTERNATIONAL, 1990, P 73-80 • C.R. HEIPLE AND J.R. ROPER, THE GEOMETRY OF GAS TUNGSTEN ARC, GAS METAL ARC, AND SUBMERGED ARC WELD BEADS, WELDING: THEORY AND PRACTICE, D.L. OLSON, R.D. DIXON, AND A.L. LIBY, ED., NORTH-HOLLAND ELSEVIER SCIENCE PUBLISHERS, AMSTERDAM, 1990, 1-34 • K.C. MILLS AND B.J. KEENE, FACTORS AFFECTING VARIABLE WELD PENETRATION,

INT.MATER.REVIEW, VOL 35, 1990, P 185-216 Transfer of Heat and Mass to the Base Metal in Gas-Metal Arc Welding Herschel B. Smartt, Idaho National Engineering Laboratory

Introduction HEAT AND MASS TRANSFER in arc welding is normally studied from the standpoint of the weld pool and heataffected zone (HAZ); however, it is also instructive to examine heat and mass transfer from the arc to the base metal. This article describes the latter topic in terms of the gas-metal arc welding (GMAW) process and provides practical information related to the development of welding procedures and the general operation of the process. Welding procedures emphasize control of parameters such as electrode speed (or current), voltage, welding speed, contact tube-to-base metal distance, as well as current pulse parameters for out-of-position welding. It is therefore easy to overlook the fact that the process is simply a source of heat and mass inputs to the weldment. Melting of the base metal, dilution of the filler metal, solidification of the weld bead, microstructural development in the weld bead and HAZ, and thermomechanical distortion and residual stresses all follow from the heat and mass inputs. The conventional parameters identified above arc variables that control the heat and mass inputs. An example of the relationship between the conventional parameters of electrode speed and welding speed to heat and mass transferred to the weldment is shown in Fig. 1 for certain conditions, as applied to the welding of thick-section steel.

FIG. 1 PLOT OF WELDING SPEED VERSUS ELECTRODE SPEED AS FUNCTION OF HEAT TRANSFER PER LENGTH OF WELD, H, AND MASS TRANSFER EXPRESSED IN TERMS OF REINFORCEMENT, G. POWER SUPPLY OPENCIRCUIT VOLTAGE, E0, IS 32 V; CONTACT TUBE-TO-BASE METAL DISTANCE, CT, IS 15.9 MM (0.625 IN.). SHADED AREA DENOTES REGION IN WHICH SPRAY AND STREAMING TRANSFER MODES OCCUR; GLOBULAR TRANSFER OCCURS AT LOWER ELECTRODE SPEEDS, AND ELECTRODE CONTACTS THE WELD POOL AT HIGHER ELECTRODE SPEEDS. SOURCE: REF 1

The issues described in this article include the: • •

TOTAL HEAT TRANSFERRED TO THE BASE METAL PARTITIONING OF HEAT TRANSFER BETWEEN THE ARC AND THE MOLTEN ELECTRODE

• • •

DROPLETS TRANSFER MODES OF THE DROPLETS ROLE OF THE ARC IN DROPLET TRANSFER SIMPLE MODEL FOR WELDING PROCEDURE DEVELOPMENT BASED UNDERSTANDING OF HEAT AND MASS TRANSFER TO THE BASE METAL

ON

AN

Reference

1. H.B. SMARTT, IDAHO NATIONAL ENGINEERING LABORATORY, 1992 Transfer of Heat and Mass to the Base Metal in Gas-Metal Arc Welding Herschel B. Smartt, Idaho National Engineering Laboratory

Heat Transfer The total transfer of heat, H, (neglecting preheating) from the GMAW process to the weldment per unit time is given

by:

H = ηEI

(EQ 1)

where E is voltage, I is current, and η is the heat-transfer efficiency. The rate at which heat is transferred to the weldment per unit length of weld is given by: H=

η EI R

(EQ 2)

where R is welding speed. Calorimeter-based heat-transfer experiments reveal that the heat-transfer efficiency for welding thick-section steel is nominally 80 to 90%, as indicated in Fig. 2. The total heat-transfer efficiency is altered somewhat by changing other parameters. For example, it increases slightly as the power supply open-circuit voltage is decreased (for a silicon controlled rectifier regulated power supply and it increases slightly with increasing contact tube-to-base metal distance. However, 85% is a reasonable estimate for most conditions.

FIG. 2 PLOT OF HEAT-TRANSFER EFFICIENCY TO BASE METAL VERSUS ELECTRODE-SPEED FOR 0.89 MM (0.035 IN.) DIAMETER STEEL ELECTRODE IN AN AR-2% O2 SHIELD GAS. TOTAL HEAT-TRANSFER EFFICIENCY

IS SHOWN PARTITIONED INTO ARC AND MOLTEN DROP COMPONENTS. POWER SUPPLY OPEN-CIRCUIT VOLTAGE, E0, IS 32 V; CONTACT TUBE-TO-BASE METAL DISTANCE, CT, IS 15.9 MM (0.625 IN.). SOURCE: REF 2

Partitioning of Heat Transfer. In the GMAW process, the molten droplets of electrode material carry a significant portion of the total heat transferred to the weld pool. This is seen in calorimetry experiments (Fig. 2), where the total heattransfer efficiency of the GMAW process is partitioned into those portions associated with transfer by the arc and by the molten droplets. At low electrode speeds, about 60% of the total heat transferred is associated with the arc. As electrode speed increases, the fraction of total heat transferred associated with the droplets increases, reaching nominally 50% at current levels in excess of about 220 A (that is, at about 230 mm/s, or 9.1 in./s, electrode speed) for the conditions used.

Reference cited in this section

2. A.D. WATKINS, "HEAT TRANSFER EFFICIENCY IN GAS METAL ARC WELDING," MASTER'S THESIS, UNIVERSITY OF IDAHO, APRIL 1989 Transfer of Heat and Mass to the Base Metal in Gas-Metal Arc Welding Herschel B. Smartt, Idaho National Engineering Laboratory

Mass Transfer Droplet Transfer Modes. Although the International Institute of Welding (IIW) lists eight distinct metal transfer modes

(Ref 3), the modes commonly used in U.S. welding practice are globular, spray, streaming, rotating, and short circuiting. The mode terms "drop" and "repelled" used by the IIW are often referred to as "globular," and the mode term "projected" is generally referred to as "spray." The globular, spray, streaming, and short-circuiting transfer modes are shown in Fig. 3, 4, 5, 6.

FIG. 3 GLOBULAR TRANSFER MODE IN GAS-METAL ARC WELDING OF STEEL. (A) SCHEMATIC SHOWING TRANSFER OF ELECTRODE MATERIAL GLOBULES ONTO CATHODE BASE METAL. (B) HIGH-SPEED PHOTOGRAPH OF GLOBULAR METAL TRANSFER

FIG. 4 SPRAY TRANSFER MODE IN GAS-METAL ARC WELDING OF STEEL. (A) SCHEMATIC SHOWING TRANSFER OF ELECTRODE MATERIAL DROPLETS ONTO CATHODE BASE METAL. (B) HIGH-SPEED PHOTOGRAPH OF SPRAY METAL TRANSFER MODE

FIG. 5 HIGH-SPEED PHOTOGRAPH OF STREAMING TRANSFER MODE IN GAS-METAL ARC WELDING OF STEEL

FIG. 6 SHORT-CIRCUITING TRANSFER MODE IN GAS-METAL ARC WELDING OF STEEL. (A) SCHEMATIC SHOWING TRANSFER OF ELECTRODE MATERIAL BY SURFACE TENSION OF WELD POOL ONTO CATHODE BASE METAL. (B) HIGH-SPEED PHOTOGRAPH OF MATERIAL TRANSFER WHEN ARC LENGTH IS VERY SHORT. (C) HIGH-SPEED PHOTOGRAPH OF CYCLIC SHORTING OF ARC BY THE ELECTRODE DURING METAL TRANSFER TO WELD POOL. (D) HIGH-SPEED PHOTOGRAPH OF VIOLENT ARC REIGNITION WITH ASSOCIATED SPATTER

When all other parameters are held constant, the metal transfer mode at the lowest wire feed speed (and associated current level) is globular. As wire speed (and therefore current) increases, the mode changes rapidly from globular to spray. With an additional increase in wire speed (and current), spray transfer becomes streaming transfer. Figure 7 shows the droplet sizes during the transition from globular to spray to streaming transfer for direct and pulsed current welding. If an adequately high current, contact tube-to-base metal distance, and voltage exist, then rotating transfer, wherein the lower part of the electrode becomes molten over a considerable length and rotates in a helical spiral under the influence of the magnetic field surrounding the arc, can occur. As it rotates, a controlled stream of droplets is transferred from the electrode tube to the weld pool over a relatively wide area. Additional increases in wire feed/current at low voltage shorten the arc length and, eventually, the wire stubs into the weld pool. In addition, with appropriate conditions (for example, carbon dioxide, argon-carbon dioxide mixtures, and helium-based shielding gases), droplets can be transferred directly, by surface tension forces, after contact of the drop with the weld pool, a condition called short-circuiting transfer.

FIG. 7 DROPLET SIZES ATTAINABLE UNDER SELECTED CURRENT PULSING CONDITIONS IN GAS-METAL ARC WELDING OF STEEL. SOURCE: REF 4

In practical applications, the optimum transfer mode depends in part on the thickness of the base metal being welded. For example, very thin sections (in all positions) require the short-circuiting mode (with low current levels and appropriate settings of voltage and other operating parameters, including shielding gas composition). Thicker sections show best results with spray or streaming transfer. These transfer modes also produce high heat input, maximum penetration, and a high deposition rate. In welding steel, they are generally limited to welding that occurs in the flat position and the horizontal fillet position, except when pulsed current is used. Rotating transfer can be used in a deep groove in thicksection material. One- and two-pass heavy fillet welds are also a major area of application for this process variation. Globular transfer (Fig. 3) involves a droplet that, generally, is much larger in diameter than the electrode wire. Globular transfer can involve a transfer rate of about 1 to 10 drops/s, and the arc has a soft, rounded appearance. Droplet detachment and transfer are mainly due to the gravitational force, which limits globular transfer to in-position welding. For spray transfer (Fig. 4), droplet and electrode diameters are roughly equivalent. Spray transfer can involve 100 drops/s, and the arc noticeably contracts, or "stiffens." Drops usually travel in-line down the center of the arc, but several drops may be in flight at the same instant. In streaming transfer (Fig. 5), a well-developed liquid column extends from the solid electrode down into the arc and breaks into small droplets before contacting the weld pool (see the discussion on spatter below). Streaming transfer can involve 1000 drops/s, and the arc has a characteristic "cone" shape, as shown in Fig. 5. In short-circuiting transfer, the arc is very short (Fig. 6a). During metal transfer, the undetached molten droplet contacts the weld pool, shorting out the arc (Fig. 6b), which then extinguishes. Surface tension plays an important role in transferring the drop to the weld pool, but detachment of the drop from the electrode is due to electromagnetic pinch forces. Reignition of the arc is violent (Fig. 6c), resulting in considerable scatter. As seen in Fig. 7, the transition from one mode to another actually involves continuous variations in droplet size. Because the product of droplet volume and transfer rate equals the electrode melting rate (in appropriate dimensions), there is also a continuous variation in droplet transfer rate. Detailed experimental studies have shown that rapid, cyclic transition from globular to spray to globular, and so on, occurs in the transition region between globular and spray transfer (Ref 5).

Although it is generally not possible for the welder to see droplet transfer events in the spray and streaming transfer modes (except under ideal conditions for spray transfer), the arc shape changes and associated changes in electrical and acoustic noise allow a trained welder to readily identify all of the transfer modes. The section "Electrical and Acoustic Signals" in this article describes this event more fully. Weld Reinforcement. Assuming that mass is conserved as welding occurs, the transverse cross-sectional area of the

weld bead added to the weldment for a single weld pass, which is called the reinforcement, G, is given by: G=

S πd² R 4

(EQ 3)

where S is electrode speed, R is welding speed, and d is electrode diameter. For the spray and streaming transfer modes, the assumption that mass is conserved is reasonably good. However, because short circuiting and other conditions generate considerably spatter, the actual reinforcement will be slightly less than the value calculated by Eq 3. Droplet Velocity and Temperature. In spray and streaming transfer, the droplets are accelerated rapidly through the

arc to the weld pool. Velocities of 1 m/s (40 in./s) are typical, increasing with voltage (Fig. 8). Calorimetry-based experimental results indicate that steel droplets can reach temperatures of approximately 2600 K (4220 °F) (Ref 2). This explains why the droplets transport one-half of the total heat transferred to the base metal.

FIG. 8 DROPLET VELOCITIES FOR SPRAY TRANSFER IN GAS-METAL ARC WELDING OF STEEL AT THREE DIFFERENT OPEN-CIRCUIT VOLTAGES. SOURCE: REF 5

A result of the activity of molten electrode droplets in heat transfer to the base metal is that they also play an important role in convective heat transport in the weld pool and, thus, in weld penetration (Ref 6). This can be seen in Fig. 9, a transverse cross section of a gas-metal arc bead-on-plate weld on carbon steel. The region of deep penetration in the center of the weld bead is associated with the heat convected to the lower portion of the weld pool by the entering droplets.

FIG. 9 TRANSVERSE CROSS SECTION OF GAS-METAL ARC BEAD-ON-PLATE WELD IN CARBON STEEL TO SHOW DEEP PENETRATION IN THE WELD BEAD CENTER GENERATED BY MOLTEN ELECTRODE DROPLETS

Electrical and Acoustic Signals. As the droplets of molten metal detach from the electrode, an almost instantaneous change occurs in the electrode extension, which results in a sudden, although small, change in the electrical resistance between the contact tube and the base metal. The results are spikes in the secondary circuit voltage and welding current, accompanied by pressure changes in the arc. Thus, both electrical and acoustic noise (Fig. 10) are generated by the GMAW process and are characteristic of the droplet transfer mode. The waveforms for globular transfer show prominent spikes associated with individual droplet transfer events. This can also be seen in Fig. 11, a plot of the computer-digitized current for a weld on carbon steel, where a change from globular to spray transfer has occurred at about 8 s into the weld. The power spectra of the current changes dramatically for the two transfer modes (Fig. 12), leading to a means of detecting transfer mode during welding.

FIG. 10 AUDIO, CURRENT, AND VOLTAGE DURING GLOBULAR TRANSFER IN GAS-METAL ARC WELDING OF STEEL. SOURCE: REF 7

FIG. 11 PLOT OF DIGITIZED CURRENT VERSUS TIME DURING A GAS-METAL ARC WELD IN CARBON STEEL IN WHICH A TRANSITION FROM GLOBULAR TO SPRAY TRANSFER OCCURRED (AT ~T = 8 S) AS CURRENT WAS INCREASED. SOURCE: REF 1

FIG. 12 WELDING CURRENT POWER SPECTRA FOR GLOBULAR TRANSFER AND SPRAY TRANSFER MODES DURING GAS-METAL ARC WELDING OF STEEL. SOURCE: REF 1

Several specific techniques have been identified (Ref 8) for detecting the mode of metal transfer: Fourier transform, standard deviation, peak ratios, and integrated amplitude of the current and voltage signal. In a similar manner, spatter, lack of shielding gas, and contact tube wear can also be detected. Similar power spectra changes also occur in the secondary circuit voltage and acoustic noise. The change in acoustic noise allows the aural detection of the transfer mode. An experienced welder can readily hear, for example, the difference between globular and spray transfer. Contact of the molten droplet with the weld pool, while the droplet is still attached to the electrode, also results in a sudden change in electrical resistance. This is important, because spatter is generally produced when this electrode-to-

weld pool contact breaks. (Spatter can also be produced by rapid expansion of gas bubbles in the electrode as it melts.) The resulting electrical and acoustic noise can be used to detect that spatter is being produced.


1. H.B. SMARTT, IDAHO NATIONAL ENGINEERING LABORATORY, 1992 2. A.D. WATKINS, "HEAT TRANSFER EFFICIENCY IN GAS METAL ARC WELDING," MASTER'S THESIS, UNIVERSITY OF IDAHO, APRIL 1989 3. J.F. LANCASTER, THE PHYSICS OF WELDING, 2ND ED., INTERNATIONAL INSTITUTE OF WELDING, PERGAMON PRESS, 1986, P 233 4. T.W. EAGAR, D.E. HARDT, H.B. SMARTT, AND J.A. JOHNSON, PROCESS CONTROLLABILITY IN GAS METAL ARC WELDING: MICRO AND MACROSCOPIC PROBLEMS, PROC. EIGHTH SYMPOSIUM ON ENERGY SCIENCES, CONF-9005183, U.S. DEPARTMENT OF ENERGY, ARGONNE NATIONAL LABORATORY, MAY 1990, P 51-67 5. D.E. CLARK, C. BUHRMASTER, AND H.B. SMARTT, DROPLET TRANSFER MECHANISMS IN GMAW, PROC. 2ND INTERNATIONAL CONFERENCE ON TRENDS IN WELDING RESEARCH (GATLINBURG, TN), 1989 6. W.G. ESSERS AND R. WALTER, HEAT TRANSFER AND PENETRATION MECHANISMS WITH GMA AND PLASMA-GMA WELDING, WELD.J., VOL 60 (NO. 2), 1981, P 37-S TO 42-S 7. J.A. JOHNSON, N.M. CARLSON, AND H.B. SMARTT, DETECTION OF METAL TRANSFER MODE IN GMAW, PROC. 2ND INTERNATIONAL CONFERENCE ON TRENDS IN WELDING RESEARCH (GATLINBURG, TN), 1989 8. G. ADAM AND T.A. SIEWERT, SENSING OF GMAW DROPLET TRANSFER MODES USING AN ER100S-1 ELECTRODE, WELD.J., VOL 69 (NO. 3), MARCH 1990, P 103-S TO 108-S Transfer of Heat and Mass to the Base Metal in Gas-Metal Arc Welding Herschel B. Smartt, Idaho National Engineering Laboratory

Procedure Development The preceding discussion leads to a logical approach toward selecting welding parameters during the development of welding procedures. First, given a nominal power supply open-circuit voltage and contact tube-to-base metal distance, the electrode speed can be set to determine the welding current. It can be shown that: S=

4η GI ( E0 + HI ) Hπ d ²

(EQ 4)

where E0 is the power supply open-circuit voltage. It is assumed that the proper selection of electrode type and shielding gas have been made, and that the electrode diameter is appropriate for the application. Second, the open-circuit voltage is adjusted to give either an acceptable arc length or electrode extension. The objective is to prevent either burnback (transfer of the arc to the contact tube) or spatter caused by shorting of the molten electrode droplets with the weld pool prior to detachment from the electrode. Fine adjustment of voltage may be necessary to obtain a clean start of the process, thus avoiding a short period of spatter or globular transfer following arc ignition. The combination of current and voltage for a given contact tube-to-base metal distance determines the droplet transfer mode. Third, the welding current determines the melting rate of the electrode. The electrode melting rate (Mrp) is a quadratic function of current (Ref 9):

M rp = (C1 + C2 A) I + C3

LI ² A²

(EQ 5)

where A is the cross-sectional area of the electrode and C1, C2, C3, and i are constants. Another researcher (Ref 10) defines Mrp by: M rp = C4 I + C5

LI ² A

(EQ 6)

where C4 and C5 are constants. Given a desired melting rate, the welding travel speed is set to obtain the desired weld bead reinforcement. The current, voltage, and weld travel speed have all now been set, thus determining the heat input per length of weld (H = ηEI/R). It should be noted that it is actually possible to independently vary heat input and mass input (in terms of weld bead reinforcement) to the weld, over at least a small range. This is shown in Fig. 13, where the same data presented in Fig. 1 are replotted in terms of reinforcement as a function of heat input per length of weld.

FIG. 13 PLOT OF REINFORCEMENT VERSUS HEAT INPUT TO SHOW PARAMETERS THAT FAVOR SPRAY TRANSFER MODE FOR GAS-METAL ARC WELDING OF STEEL. POWER SUPPLY OPEN-CIRCUIT VOLTAGE, E0, IS 32 V; CONTACT TUBE-TO-BASE METAL DISTANCE, CT, IS 15.9 MM (0.625 IN.).

The above discussion on procedure development ignores second-order effects, such as the dependence of arc length on weld travel speed. It should also be realized that several iterations through the steps defined above may be required for final parameter determination. Proper power supply settings must be used, and code requirements must be met.


9. A. LESNEWICH, CONTROL OF MELTING RATE AND METAL TRANSFER IN GAS-SHIELDED

METAL-ARC WELDING, PART II--CONTROL OF METAL TRANSFER, WELD.J., VOL 37 (NO. 9), SEPT 1958, P 418-S TO 425-S 10. J.H. WASZINK AND G.J.P.M. VAN DEN HEUVEL, HEAT GENERATION AND HEAT FLOW IN THE FILLER METAL IN GMA WELDING, WELD.J., VOL 61 (NO. 8), AUG 1982, P 269-S TO 282-S Transfer of Heat and Mass to the Base Metal in Gas-Metal Arc Welding Herschel B. Smartt, Idaho National Engineering Laboratory

References

1. H.B. SMARTT, IDAHO NATIONAL ENGINEERING LABORATORY, 1992 2. A.D. WATKINS, "HEAT TRANSFER EFFICIENCY IN GAS METAL ARC WELDING," MASTER'S THESIS, UNIVERSITY OF IDAHO, APRIL 1989 3. J.F. LANCASTER, THE PHYSICS OF WELDING, 2ND ED., INTERNATIONAL INSTITUTE OF WELDING, PERGAMON PRESS, 1986, P 233 4. T.W. EAGAR, D.E. HARDT, H.B. SMARTT, AND J.A. JOHNSON, PROCESS CONTROLLABILITY IN GAS METAL ARC WELDING: MICRO AND MACROSCOPIC PROBLEMS, PROC. EIGHTH SYMPOSIUM ON ENERGY SCIENCES, CONF-9005183, U.S. DEPARTMENT OF ENERGY, ARGONNE NATIONAL LABORATORY, MAY 1990, P 51-67 5. D.E. CLARK, C. BUHRMASTER, AND H.B. SMARTT, DROPLET TRANSFER MECHANISMS IN GMAW, PROC. 2ND INTERNATIONAL CONFERENCE ON TRENDS IN WELDING RESEARCH (GATLINBURG, TN), 1989 6. W.G. ESSERS AND R. WALTER, HEAT TRANSFER AND PENETRATION MECHANISMS WITH GMA AND PLASMA-GMA WELDING, WELD.J., VOL 60 (NO. 2), 1981, P 37-S TO 42-S 7. J.A. JOHNSON, N.M. CARLSON, AND H.B. SMARTT, DETECTION OF METAL TRANSFER MODE IN GMAW, PROC. 2ND INTERNATIONAL CONFERENCE ON TRENDS IN WELDING RESEARCH (GATLINBURG, TN), 1989 8. G. ADAM AND T.A. SIEWERT, SENSING OF GMAW DROPLET TRANSFER MODES USING AN ER100S-1 ELECTRODE, WELD.J., VOL 69 (NO. 3), MARCH 1990, P 103-S TO 108-S 9. A. LESNEWICH, CONTROL OF MELTING RATE AND METAL TRANSFER IN GAS-SHIELDED METAL-ARC WELDING, PART II--CONTROL OF METAL TRANSFER, WELD.J., VOL 37 (NO. 9), SEPT 1958, P 418-S TO 425-S 10. J.H. WASZINK AND G.J.P.M. VAN DEN HEUVEL, HEAT GENERATION AND HEAT FLOW IN THE FILLER METAL IN GMA WELDING, WELD.J., VOL 61 (NO. 8), AUG 1982, P 269-S TO 282-S Arc Physics of Gas-Tungsten Arc Welding J.F. Key, EG&G Idaho, Inc.

Introduction THE GAS-TUNGSTEN ARC WELDING (GTAW) process is performed using a welding arc between a nonconsumable tungsten-base electrode and the workpieces to be joined. C.E. Jackson defined a welding arc as "a sustained electrical discharge through a high-temperature conducting plasma producing sufficient thermal energy so as to be useful for the joining of metals by fusion." This definition is a good foundation for the discussion that follows.

The physics of GTAW are fundamental to all arc processes and are more straightforward, because the complications of materials (for example, filler and flux) transferred through and interacting with the arc can be avoided. Geometrically, the arc discharge in GTAW is between a rod-shaped tungsten electrode and a planar-shaped electrode, that is, the workpiece. Pure tungsten electrodes are less expensive and, possibly, more environmentally compatible than those with rare earth or other oxide additions. They are used for lower-specification welds, where tungsten contamination that is caused by the molten electrode surface can be tolerated. They are also used for alternating current (ac) welding of aluminum, copper, magnesium, and thin sections of low-alloy and stainless steels. Analysis of the arc discharge is separated into electrode regions and the arc column. The electrode regions are confined to very small distances from the electrode surfaces, have very high electrical and thermal fields, and have much higher current density, because of the contraction of the arc to a small spot. As a result, electrode regions for both the cathode and the anode are difficult to analyze by diagnostic measurements and theoretical computation. This situation must be remedied for a thorough understanding of the process, because the process parameters control the arc discharge at the cathode, with the anode serving as the connection to ground. The arc column, on the other hand, is relatively easy to analyze, but is important primarily as a means to deduce arc characteristics at the electrodes. Polarity. The GTAW process generally utilizes a direct current (dc) arc, where the tungsten electrode has a negative polarity. The tungsten electrode thus becomes the cathode and the workpiece becomes the anode. The polarity is called straight polarity, or direct current electrode negative (DCEN).

Reverse polarity, or direct current electrode positive (DCEP), is literally the reverse of DCEN. The workpiece is the cathode and the tungsten electrode serves as the anode. Because most heat is generated at the anode in the GTAW process, DCEP is used for welding certain thin-section, low melting point materials when DCEN would be likely to cause excessive penetration or burn-through. Either alternating current or DCEP is used for removing an oxide film from the surface of the weld pool or workpiece. The oxide film promotes emission during the half-cycle (ac) when the workpiece is negative polarity. As the oxide is depleted, the emission moves to a new location that has a high enough oxide content to sustain the discharge of electrons. The arc root or cathode spot where the emission occurs is highly mobile in ac or DCEP and, as a result, the arc is much less stable than in DCEN. Gas Shielding. In all cases, the arc and both electrodes are shielding by gas, usually an inert gas or a gas mixture. Argon

and argon-helium mixtures are used most often, although argon-hydrogen mixtures are used for some applications. The GTAW process may simply utilize the arc to fuse the workpieces together without the addition of filler materials (autogenous) or filler may be added to the molten pool to fill grooves in thicker weldments. A reasonable understanding of welding arc fundamentals and the GTAW process requires a more thorough discussion of the electrode regions of the arc and the arc column. Arc Physics of Gas-Tungsten Arc Welding J.F. Key, EG&G Idaho, Inc.

Electrode Regions and Arc Column The cathode and anode are similar in several respects. Both exhibit a voltage drop caused by a space charge that covers a very thin region over their surfaces, and the arc is significantly contracted on the surfaces. Figure 1 shows that the total arc voltage is partitioned between the electrode drops and arc column. The relative magnitude of these drops depends on welding parameters and electrode material.

FIG. 1 PLOT OF RELATIVE ARC VOLTAGE DISTRIBUTION VERSUS RELATIVE ARC LENGTH BETWEEN ELECTRODES

The arc discharge requires a flow of electrons from the cathode through the arc column to the anode, regardless of polarity or whether ac or dc is used. Two cases of electron discharge at the cathode will be discussed: thermionic emission and nonthermionic emission, also called cold cathode, or field emission. Thermionic emission results from joule heating (resistance) of the cathode by the imposed welding current until the electron energy at the cathode tip exceeds the work function (energy required to strip off an electron). This case applies to the general case of DCEN, where the tungsten electrode is the emitter, or cathode. Pure tungsten electrodes have to be heated to their melting point to achieve thermionic emission. Once molten, the equilibrium tip shape becomes a hemisphere, and a stable arc results from uniform emission over this surface. Thoria (ThO2), zirconia (ZrO2), or ceria (CeO2) are added to pure tungsten in amounts up to 2.2 wt% ThO2, 0.4 wt% ZrO2, or 3.0 wt% CeO2 to lower the work function, which results in thermionic emission at lower temperatures and avoids melting the cathode tip. These electrodes typically have a ground conical tip, and thermionic emission is localized to a cathode spot. Thermionic emission creates a cloud of electrons, called a space charge, around the cathode. If a second electrode at a higher potential is nearby (the workpiece, in this case), then the electrons will flow to it, thus establishing the arc. Nonthermionic, or field, emission creates an electron discharge with a very high electric field, typically exceeding 109 V/m. This intense electric field literary pulls electrons out of a relatively cold or unheated cathode. This would not appear to be applicable to welding until one considers that for reverse polarity or DCEP, a condensation of positive ions from the arc column can build up in a very thin (1 nm, or 0.04 μin.) layer over the cathode surface, creating a very high localized electric field even though the cathode voltage drop may only amount to several volts. An oxide layer, which is always present on the cathode surface in an actual weld, facilitates the discharge with a source of lower work function electrons. When the oxide layer is very thin (on the order of one atom layer), emission occurs via a tunneling mechanism through the film to an emitting site. Thicker oxide films exhibit locally conducting spots at the end of filamentary channels through the oxide. Large currents flow in these channels, which are on the order of 100 nm (4 μin.) in diameter and have lifetimes of 0.001 to 1 μs. The cathode cleaning action, which is one of the principal reasons to use DCEP or ac, results from stripping away the oxide film at the emitting sites by very small and intense jets of metal vapor and debris. It becomes obvious that a practical implication of the short lifetime of these cathode spots is a generally unstable arc that is due to the necessity of continual movement of the cathode spot to undepleted regions of oxide film. Arc instability is undesirable and DCEP or ac is only used when cathodic cleaning or the minimizing of heat input to the workpiece is a higher priority than optimizing weld bead shape and location. Anode. Welding process parameters (for example, current and voltage) control the arc discharge at the cathode. Although

the electron flow enters the anode through the anode spot and constitutes 85% of the energy going into the weld pool, thus making current density the single most important welding parameter that determines pool shape, events at the anode can

only be controlled indirectly by controlling the cathode. Anode spot stability does depend somewhat on shielding gas composition and the shape of the anode (that is, weld groove). Current density and heat input measurements at the anode have been made to better understand how process parameters that are largely controlling events at the cathode will, in turn, influence the shape and melting rate of the weld pool. The relative contributions of heat transfer to the workpiece, in terms of the GTAW process, are shown in Fig. 2.

FIG. 2 RELATIVE HEAT TRANSFER CONTRIBUTIONS TO WORKPIECE WITH GTAW. (A) CONTRIBUTION OF INDIVIDUAL PARAMETERS TO ANODE HEAT INPUT. (B) HEAT OUTPUT AT CATHODE (WORKPIECE) RELATIVE TO WELD POOL HEAT LOSS

The Thomson effect represents the energy lost by electrons as they move from higher to lower temperatures. The sum of work function, Thomson effect, and anode fall gives an electron contribution to heat transfer of approximately 84%. The remaining 16% is due to thermal effects (that is, conduction, convection, and radiation). There are small heat losses from the pool that are due to evaporation of metal ions and radiation. Figure 3 shows that the energy distribution for this particular case approaches a Gaussian distribution (that is, normal distribution curve). High helium contents in the shielding gas have produced data that are typically better fit by a Lorentzian curve, indicating a narrower current density distribution (that is, a more-contracted arc at the anode spot).

FIG. 3 PLOT OF ELECTRON AND THERMAL CONTRIBUTIONS TO HEAT TRANSFER. A, TOTAL ARC POWER (STANDARD DEVIATION, σ , OF 0.8 MM, OR 0.031 IN.); B, ELECTRON CONTRIBUTION (σ = 0.7 MM OR 0.028 IN.); C, THERMAL CONTRIBUTION. WELD PARAMETERS: CURRENT, 10 A; VOLTAGE, 10 V; TIME, 10 S; SHIELDING GAS, ARGON; ELECTRODE ANGLE, 30°

Arc efficiency, in addition to those variables that have an effect on it, is an extremely important term in the heat transfer analysis of welding. It gives the percentage of heat dissipated in the arc that actually is captured by the workpieces and is available for melting. Arc efficiency, as a function of all GTAW welding parameters and many materials, has been determined experimentally and found to be nominally 75%. The variables having the greatest effect on arc efficiency are arc voltage and anode material. For those variables, the effect is usually no more than ±5%. Other parameters have a negligible effect. Arc Column. The electron discharge between the electrodes partially ionizes the shielding gas in its path, thus making

the arc column a conductor, or plasma. Overall, the arc column is neutral and is composed of electrons, positive gas and possibly metal ions, and neutral gas atoms. Ironically, fundamental measurements of arc properties are most easily made in the arc column, although the actual effect of these properties on the electrode region of an actual weld must still be inferred. Nevertheless, it is useful to understand fundamentals of the arc that relate essential welding variables (for example, current, voltage, electrode gap, choice of shielding gas, and electrode shape) to arc temperature, current density distribution, and gas flow structure at the anode surface. Effect of Cathode Tip Shape. For the general case of straight polarity, DCEN, the tip of the tungsten alloy cathode is

ground to a point and then truncated somewhat to prevent the sharp tip from burning off and contaminating the weld. The included angle of the cone and the diameter of the truncation under some welding conditions have a significant effect on weld pool shape. Figure 4 shows examples of the effect of these two parameters on weld pool shape.

FIG. 4 FUSION ZONE PROFILE FOR SPOT-ON-PLATE WELDS AS A FUNCTION OF ELECTRODE TIP GEOMETRY USING 100% AR AS A SHIELDING GAS. WELD PARAMETERS: CURRENT, 150 A; DURATION, 2 S

For a stationary spot-on-plate weld shielded by pure argon, the weld depth-to-width ratio increased with an increasing vertex angle up to 90° and with an increasing truncation diameter. The arc became less "bell shaped" and more "ball shaped" as the vertex angle or truncation diameter increased. These results should be a valid indication of the effect of cathode tip shape for pulsed current welding, which produces a series of overlapping spot welds. A study of bead-on-plate welds (Fig. 5) made with constant current and velocity indicated a similar but less pronounced trend. These conditions produce a tear-drop molten pool shape when viewed from above, compared to circular shape for spot or pulsed current welding. Fluid and heat flow within the pool is less uniform front-to-rear in a tear-drop-shaped pool and probably has a greater influence on pool shape than electrode tip shape.

FIG. 5 FUSION ZONE PROFILE FOR BEAD-ON-PLATE WELDS AS A FUNCTION OF ELECTRODE TIP GEOMETRY USING 100% AR AS A SHIELDING GAS. WELD PARAMETERS: CURRENT, 150 A; WELDING SPEED, 3 MM/S (0.12 IN./S)

When the arc is used in a weld groove, the relative shapes of the cathode tip and the anode groove become more important. The arc discharge from the cathode will seek a path to ground with the lowest electrical resistance. For a stable arc properly centered in the groove (for example, a root pass), the shortest path to ground should be between the cathode tip and the bottom of the groove (Fig. 6). This will require the cathode vertex angle to be somewhat less than the included angle of the groove and/or a sufficiently wide groove to ensure that the shortest path to ground is from the cathode tip to the groove bottom and not, for example, from the electrode shoulder to the groove wall, as the case would be with a 90° electrode in a 10° narrow groove. Welding in a groove places a higher priority on arc stability and location than on maximum penetration.

FIG. 6 EFFECTS OF ELECTRODE TIP GEOMETRY ON THE PATH LENGTH TO GROUND IN WELD GROOVES OF VARIOUS SHAPES. (A) 75° V-GROOVE. (B) 40° U-GROOVE. (C) 10° NARROW GROOVE

To understand the effects of tip shape, temperature distributions in the plasma were measured. Figure 7 shows that as the cathode vertex angle increases, the plasma radius of the arc column increases at mid-gap and becomes more constricted near the anode. The quantitative interpretation of these results requires theoretical modeling, which has yet to be completed.

FIG. 7 EFFECT OF VERTEX ANGLE ON GTAW ARC COLUMN TEMPERATURE DISTRIBUTION WITH 100% AR USED AS SHIELDING GAS. (A) 30° ELECTRODE VERTEX ANGLE. (B) 90° ELECTRODE VERTEX ANGLE. WELDING CURRENT, 150 A

Effect of Shielding Gas Composition. The GTAW process typically uses either an inert gas, such as argon, or an inert gas mixture, such as argon and helium, to shield the arc and the weld from atmospheric contamination. Occasionally, a slightly reactive gas mixture, such as argon with up to 15 vol% H2, is used. (The 15 vol% limit is based on safety considerations.) Shielding gas composition has a rather strong effect on arc temperature distribution and, under certain conditions, a significant effect on weld pool shape.

Figure 8 shows how shielding gas affects arc voltage. The curves would all be displaced downward for shorter arc lengths, but the relative positions would be maintained. Figure 9 shows the effects of both cathode vertex angle and shielding gas composition on weld pool shape for spot-on-plate welds. Increasing additions of helium to argon show a remarkable increase in penetration when using a 30° vertex angle. However, the effect is much less evident when using a 90° vertex angle.

FIG. 8 PLOT OF ARC VOLTAGE VERSUS ARC CURRENT FOR SELECTED INERT SHIELDING GASES. WELDING PARAMETERS: ANODE, TITANIUM; CATHODE, TUNGSTEN; POLARITY, DCEN; ARC LENGTH, 12.7 MM (0.050 IN.)

FIG. 9 EFFECT OF ELECTRODE TIP GEOMETRY AND SHIELDING GAS COMPOSITION ON WELD POOL SHAPE FOR SPOT-ON-PLATE WELDS. WELDING PARAMETERS: CURRENT, 150 A; DURATION, 2 S

To understand these phenomena, arc temperature distributions for a variety of shielding gases and mixtures, electrode shapes, current, arc voltages (electrode gaps), and anode materials have been measured in order to clarify welding arc fundamentals. Welding arcs are composed of electrons, positive gas ions, and neutral gas atoms. Some measurement techniques give the temperature of one species (electrons), whereas other techniques give the temperature of another species (neutral atoms). If the arc is in local thermodynamic equilibrium (LTE), all techniques should give the same temperature. Although the assumption of LTE used to be considered completely valid, contemporary investigations have suggested that this is not always the case and that some of the older measurements may be somewhat in error. Absolute values of arc temperature are only needed to establish boundary conditions for detailed arc modeling of temperature-dependent properties. What is of more importance to the welding engineer or technologist is the relative effect of essential variables

on heat input to the workpiece. Arc temperature measurement is one useful indication of how these variables affect the arc. When compared to Fig. 7(a), Fig. 10 shows that large additions of helium to argon decrease peak temperatures slightly and increase the plasma diameter in the plasma column. The arc appears to become a broader and more isothermal heat source. The lower peak temperature is reasonable, because a combination of the high ionization potential of helium and relatively low currents of welding arcs gives an arc column that is only slightly ionized. Figure 7 showed that a large vertex angle had a similar, but less pronounced, effect on arc temperature when adding helium to the shielding gas (that is, the axial peak temperature decreased and the plasma diameter increased).

FIG. 10 PLOT OF GTAW ARC COLUMN TEMPERATURE DISTRIBUTION AS A FUNCTION OF ANODE DISTANCE AND ARC POSITION. WELDING PARAMETERS: ELECTRODE VERTEX ANGLE, 30°; CURRENT, 150 A; SHIELDING GAS, 10AR-90HE

Figure 11 shows that doubling the current from 150 to 300 A produces an increase in plasma diameter (that is, that portion of the arc above approximately 8000 K, or 13,900 °F). This effect occurs regardless of the shielding gas composition.

FIG. 11 PLOT OF GTAW ARC COLUMN TEMPERATURE DISTRIBUTION RELATIVE TO ANODE DISTANCE AND ARC POSITION. WELDING PARAMETERS: ELECTRODE VERTEX ANGLE, 30°; CURRENT, 300 A; SHIELDING GAS, 100% AR

The arc length or gap between the electrodes is yet another process parameter that must be considered, especially for mechanized welding, where it can be kept reasonably constant. Because arc length is proportional to arc voltage, longer arcs have higher arc voltages and consume more energy for the same current. However, this increased energy is generally lost through radiation to the environment surrounding the weld and does not effectively supply additional heat to the workpieces. Mechanized welding generally utilizes rather short arc lengths (2 to 3 mm, or 0.08 to 0.12 in.), whereas manual welding uses a longer arc length. Unfortunately, relationships that establish a direct correlation between the temperature distribution in the arc column and the weld pool shape still have not been established, primarily because weld pool shape depends on other factors, such as compositionally dependent molten metal properties. Flow Structure. Shielding gas is used to displace reactive gases in the atmosphere from the vicinity of the weld. Inert

gases are preferred for the GTAW process, because they minimize unwanted gas-metal reactions with the workpieces. A uniform laminar flow of gas from the gas cup would be ideal and, in fact, is usually achieved as long as there is no welding arc. However, the arc discharge rapidly heats the gas in the arc column, and thermal expansion causes plasma jets. The lower temperature near the cathode tip in Fig. 11 is an indication of a jet pumping in cooler gas. This becomes an important factor at high currents, because these jets can depress the surface of the weld pool and alter heat transfer to it. The rapid gas expansion can cause the flow to deviate from laminar and, in extreme cases, the flow can become turbulent. Turbulence tends to mix atmospheric contaminants into the arc, often where they can do the most harm: at the surface of the molten weld pool. Holographic interferometry and Schlieren photography have been used to characterize gas flow. Figures 12, 13, and 14 show examples of laminar and turbulent flow. The flow from a commercial-design gas cup for three current levels is shown in Fig. 12. Laminar flow occurs where the fringes are generally straight and parallel. Turbulent flow is indicated by very wavy or circular fringes. Increasing current tends to make the laminar region somewhat broader and thicker, effectively increasing the area shielded from atmospheric contamination.

FIG. 12 EFFECT OF GEOMETRY ON COMMERCIAL GAS CUP LAMINAR AND TURBULENT FLOW AS DETECTED BY REAL-TIME HOLOGRAPHIC INTERFEROMETRY

FIG. 13 EFFECT OF GEOMETRY ON CONVERGING CONE CUP LAMINAR AND TURBULENT FLOW AS DETECTED BY REAL-TIME HOLOGRAPHIC INTERFEROMETRY

FIG. 14 EFFECT OF GEOMETRY ON VENTURI GAS CUP LAMINAR AND TURBULENT FLOW AS DETECTED BY REAL-TIME HOLOGRAPHIC INTERFEROMETRY

Figures 13 and 14 result from experiments with gas cup shapes that were designed to improve shielding. Figure 13 shows that a converging cone would be a very poor choice for the GTAW process, as indicated by the very small area of laminar flow and the extreme turbulence in the surrounding region.

Figure 14 is a venturi shape, which provides a large laminar flow region for all current levels and excellent shielding. This design may be somewhat better than commercial designs, but weld contamination studies should be conducted to verify this possibility. Arc Physics of Gas-Tungsten Arc Welding J.F. Key, EG&G Idaho, Inc.

Selected References

• S.S. GLICKSTEIN, "ARC-WELD POOL INTERACTIONS," DEPARTMENT OF ENERGY RESEARCH AND DEVELOPMENT REPORT WAPD-TM-1429, BETTIS ATOMIC POWER LABORATORY, AUG 1978 • C.E. JACKSON, THE SCIENCE OF ARC WELDING,WELD. J., VOL 39, 1960, P 129S-140S, 177S-190S, 225S-230S • J.F. KEY, H.B. SMARTT, J.W. CHAN, AND M.E. MCILWAIN, PROCESS PARAMETER EFFECTS ON ARC PHYSICS AND HEAT FLOW IN GTAW, PROC. INT. CONF. ON WELDING TECHNOLOGY FOR ENERGY APPLICATIONS (GATLINBURG, TN), 16-19 MAY, 1982, P 179-199 • J.F. KEY, J.W. CHAN, AND M.E. MCILWAIN, PROCESS VARIABLE INFLUENCE ON ARC TEMPERATURE DISTRIBUTION, WELD. J./WELD. RES. SUPP., VOL 62 (NO. 7), P 179-S TO 184-S • J.F. LANCASTER, ED., THE PHYSICS OF WELDING, 2ND ED., INTERNATIONAL INSTITUTE OF WELDING, PERGAMON PRESS, 1986 • H.B. SMARTT, J.A. STEWART, AND C.J. EINERSON, "HEAT TRANSFER IN GAS TUNGSTEN ARC WELDING," ASM METALS/MATERIALS TECHNOLOGY SERIES, ASM INTERNATIONAL WELDING CONGRESS (TORONTO), 14-17 OCT 1985 Power Sources for Welding F. James Grist, Miller Electric Manufacturing Company, Inc.; William Farrell and Glen S. Lawrence, Ferranti Sciaky

Introduction POWER SOURCES are apparatuses that are used to supply current and voltages that are suitable for particular welding processes. This article describes power sources for arc welding, resistance welding, and electron-beam welding. Power Sources for Welding F. James Grist, Miller Electric Manufacturing Company, Inc.; William Farrell and Glen S. Lawrence, Ferranti Sciaky

Arc Welding Power Sources Arc welding requires that an electric arc be established between an electrode and the workpiece to produce the heat needed for melting the base plate. Because utility energy is not delivered at the proper voltage and current, it must be converted to the required levels by the welding power source. Arc power sources convert the customary 240 or 480 V alternating current (ac) utility power to a range from 20 to 80 V and simultaneously increase the current proportionately. Galvanic isolation is also provided from the utility line voltage to ensure operator safety. Motor- or engine-driven welding generators are wound to deliver the correct voltage and current

directly; therefore, no transformer is necessary. Actual arc voltage is established by factors other than the power source (for example, electrode type, base-metal type, shielding gas, and standoff distance). Ratings and Standards The National Electrical Manufacturers Association (NEMA) categorizes arc welding power sources into three classes on the basis of duty cycle: • • •

CLASS I: RATED OUTPUT AT 60, 80, OR 100% DUTY CYCLE CLASS II: RATED OUTPUT AT 30, 40, OR 50% DUTY CYCLE CLASS III: RATED OUTPUT AT 20% DUTY CYCLE

Duty cycle is the ratio of arc time to total time based on a 10 min averaging period. A 60% machine will deliver 6 min of arc time and 4 min off time without overheating. In Fig. 1, curve A shows a NEMA class I (60%) 300 A rated machine that is capable of a maximum 375 A at reduced duty cycle (38%) and 232 A at 100% (continuous). Curve B represents a NEMA class II (50%) 250 A machine with a continuous duty of 176 A. Curve C represents an engine-driven machine rated at 225 A and 20% duty. It does not offer output in excess of its rating, because of a horsepower limitation of the engine.

FIG. 1 SELECTED DUTY CYCLE CURVES. A, 300 A, 60% MACHINE; B, 250 A, 50% MACHINE; C, 225 A, 20% ENGINE-DRIVEN MACHINE WITH HORSEPOWER LIMITATION

NEMA requires most machines to produce a maximum of 125% of rated output current. No minimum current is specified, but 10% (of rating) is typical. Load voltage, E, for class I and II machines is defined by:

E = 20 + 0.04I

(EQ 1)

where I is the current in amperes, below 600 A. Above 600 A, the rated load voltage remains at 44 V. Power sources intended for gas-tungsten arc welding (GTAW) service are defined by:

EGTAW = 13 + 0.012I Class III machines are constant-current types limited to 20% duty cycle.

(EQ 2)

Efficiencies (that is, ratios of power output to power input) of 68% for line-frequency machines and 88% for inverters are typical. Line current draw can be reduced in certain machines by the addition of power factor correcting capacitors connected to the primary winding. Input (line) voltages shown on the manufacturer nameplate are often multivoltage by means of tap changing for 200, 230, 460, or 575 V. Designed primarily for 60 Hz power, most systems will operate at 50 Hz, with minor changes in output. Input voltages for 50 Hz systems are typically 220, 380, and 415 V. Usual service conditions are temperature that range from 0 to 40 °C (32 to 105 °F) and altitudes of up to 1000 m (3300 ft). (Altitude is a factor because of fuel-air mixtures.) The power source nameplate must include the following specifications: • • • • • • • • •

MANUFACTURER NEMA CLASS (I, II, OR III) AND DUTY CYCLE MAXIMUM OPEN-CIRCUIT VOLTAGE NUMBER OF PHASES MAXIMUM SPEED (IN REVOLUTIONS PER MINUTE) AT NO LOAD (ROTATING EQUIPMENT) RATED LOAD (IN VOLTS) RATED LOAD (IN AMPERES) FREQUENCY (OR FREQUENCIES) IN HERTZ INPUT VOLTAGE(S)

The National Electric Code (NEC) requires that a main disconnect switch be provided for each installed welding machine. Manufacturer installation instructions list the kVA and kW rating and recommend the correct fuse and conductor size. Local codes may differ from NEC requirements. Proper grounding is a necessary safety consideration. Machines equipped with high-frequency arc stabilizers must be correctly grounded to prevent electromagnetic interference (EMI) to nearby equipment. The installer certifies (to the Federal Communications Commission) that the machine has been grounded according to manufacturer instructions. Most power sources are forced-air cooled by internal fans. Thermal sensors will interrupt output if airflow is impeded. Power Source Selection Because no single power source is right for all welding situations, it is necessary to know the processes to be used before selecting the best power source. Table 1 gives the more-common processes that use constant-current and constant-voltage power sources.

TABLE 1 SELECTION OF POWER SOURCE RELATIVE TO WELDING PROCESS TO BE EMPLOYED

WELDING PROCESS

POWER SOURCE CONSTANT CONSTANT CURRENT VOLTAGE SHIELDED METAL ARC WELDING (SMAW) • ... GAS-TUNGSTEN ARC WELDING (GTAW) • ... GAS-METAL ARC WELDING (GMAW) ... • PULSED GAS-METAL ARC WELDING (GMAW-P) • ... FLUX-CORED ARC WELDING (FCAW) ... • SUBMERGED ARC WELDING (SAW) ... • ELECTROGAS WELDING (EGW) • • ELECTROSLAG WELDING (ESW) ... • PLASMA ARC WELDING (PAW) • ...

PLASMA ARC CUTTING (PAC)

•

...

Alternating current power sources for shielded metal arc welding (SMAW) and submerged arc welding (SAW) can be as simple as a single transformer (Fig. 2a). Low-voltage machines with open-circuit voltages of 50 V or less are transformer types. Output alternating current reflects the sinusoidal waveform of the utility input line (Fig. 3a).

FIG. 2 SELECTED ALTERNATING CURRENT POWER SOURCES. (A) ALTERNATING CURRENT WELDING TRANSFORMER, ADJUSTABLE CORE OR WINDINGS. (B) THREE-PHASE SCR-CONTROLLED DC WELDING POWER SOURCE. (C) INVERTER BLOCK DIAGRAM. (D) MOTOR-GENERATOR SET. (E) ENGINE-DRIVEN ALTERNATOR WITH DC OUTPUT. (F) SECONDARY SWITCHING POWER SOURCE

FIG. 3 TYPICAL AC WELDING CURRENT WAVEFORMS. (A) SINE WAVE ALTERNATING CURRENT, SIMPLE TRANSFORMER WELDER. (B) MODIFIED SINE WAVE OBTAINED FROM MAGNETIC AMPLIFIER-TYPE WELDER. (C) SQUARE WAVE, LINE FREQUENCY, AND BALANCED DWELL. (D) SQUARE WAVE WITH UNBALANCED DWELLS, EP VERSUS EN. (E) VARIABLE-POLARITY AC WITH DWELL, AMPLITUDE, AND FREQUENCY ALL ADJUSTABLE

Magnetically controlled equipment (that is, a saturable reactor or magnetic amplifier) used for GTAW delivers a modified sinewave output (Fig. 3b). There is a risk of electrode (tungsten) migration to the workpiece if peak current greatly exceeds the average value. This technology exhibits moderate response. Newer technologies deliver a square wave output at line frequency (Fig. 3c). A square wave eliminates peaking and provides a rapid transition through zero, which is important to cyclic reignition of the arc. Thyristors are employed in concert with magnetic cores to generate the square current waveform. Adjustable imbalance permits the operator to control the ratio of electrode positive (EP) to electrode negative (EN) current by dwell extension (Fig. 3d). Direct current power sources can be classified as being constant-current sources or constant-voltage sources. Constant-Current Sources. For the SMAW process, a family of "drooping" volt-ampere (V-A) curves (A, B, and C) is

provided in Fig. 4(a). Often, a current boost or "dig" characteristic is made available (curve D in Fig. 4a). A burst of extra current during the brief period of metal transfer shorting tends to alleviate sticking during keyhole welding or when weaving or whipping the electrode, especially with the more-globular transfer electrodes.

FIG. 4 VOLT-AMPERE CURVES OF TYPICAL DC POWER SOURCES. (A) CONSTANT-CURRENT SOURCE. (B) CONSTANT-VOLTAGE SOURCE

Power sources for dc GTAW processes are generally more critically constant-current than are the SMAW "stick droopers." Ripple content is lower and current regulation is tighter. Inverter types customarily regulate to ±1% or ±1 A, whichever is larger. Remote foot control is essential to ensure that both hands are free for manipulation. Response time need not be fast, except for pulsed programs. Programmable and pulse-capable units permit the automation of complex current profiles. Power sources for GTAW processes include a high-frequency arc starting device that impresses a high radio frequency (RF) voltage on the electrode. This energy "jumps the gap" from the electrode to the workpiece, ionizing the shielding gas, and permits establishment of an arc. Thus, the electrode need not touch the workpiece. Constant-Voltage Sources. Power sources intended for gas-metal arc welding (GMAW) exhibit a relatively flat V-A curve, as shown in Fig. 4(b) (curves E, F, and G). The self-correcting characteristic of GMAW regulates the electrode burn-off rate. As the wire feed rate is increased, additional current will be drawn, but terminal voltage will remain

essentially constant. Voltage feedback techniques are used in solid-state designs to hold long-term output close to a set value. In the short-circuiting GMAW process, for example, drops of metal are transferred to the base metal through a pinch-off routine during a period of short circuiting of the electrode to the workpiece. Shorting time and arc time are determined by power source dynamics. Source Characteristics Whether utility fed or motor driven, all welding power sources have a characteristic that indicates their dynamic and steady-state response. In early machines, source characteristics were optimized for a single process and were relatively unchangeable. More recent designs can be programmed for a variety of characteristics that include multiple weld processes. The welding arc process is a dynamic, fast-changing load, exhibiting short circuits (if shorting, then no arc) at high currents and momentary voltages from near zero to over 50 V. Changes occur in milliseconds, particularly when metal transfer takes place. The ability of the power source to respond properly to rapidly changing conditions at the arc is critical to good performance. Modern technology invariably includes a closed-loop feedback circuit in the control system. Open-Circuit Voltage (OCV). When no load is connected to the output terminals of a welding power source, the

voltage that appears at the terminals is at its maximum. A high OCV value generally aids in arc starting and stability. In transformer-type power sources. OCV is established by the incoming utility line voltage and the transformer primary-tosecondary turns ratio. The output winding turns, the flux, and the angular velocity (rev/min) will establish the OCV in generators. In manual welding, the maximum OCV is 80 Vrms for machines with more than 10% ripple, or 100 V OCV for a ripple content of less than 10%. Power sources intended for automatic welding only are permitted 100 V OCV. The percentage ripple voltage is the ratio of rms ripple to average pulsating dc, expressed as: %ripple =

Vmax − Vmin Vr m s

(EQ 3)

Static V-A Curves. Two terms are commonly used to define the characteristics of a power source for welding: static

output and dynamic response. Static characteristics are set forth by the manufacturer as a series of V-A curves (Fig. 4). The slope of the static V-A curve is the ratio of load voltage change to the change in load current, expressed in V/100 A. It indicates the achievable envelope of output limits, including the highest voltage the machine will support at a given current, and the highest current that can be drawn at a given voltage. Dynamic Response. The static V-A curves discussed above do not take into account the dynamic interplay between the

power source and the process. Instantaneous response of the power source to step changes in load voltage is important to welding performance. This is especially true in the case of certain GMAW processes, where the response to a metal transfer short circuit is instrumental in bead formation and the reduction of spatter generation. Methods of Control or Electrical Conversation. Power source control methods are a key factor in the rate of current

rise (dI/dt). Either increased reactance in the ac circuitry or inductance in the dc portion will slow the rate of rise (and fall) of current during changes at the arc. Some equipment incorporates an adjustment to select the proper number of turns for a given welding condition (with silicon-controlled rectifiers [SCR] and transistor controls, these characteristics can be synthesized). Because there is no industry-accepted method to quantify the dynamics, the user must obtain a power source with a satisfactory range of adjustment. Another consideration is the frequency response of the power source, that is, its ability to "follow" a fast command signal, such as a pulse program. Magnetically based machines are generally not fast enough to faithfully reproduce a pulsed GMAW control signal and deliver it at high currents to the weld. When a better response is required, inverter-based SCR, transistorized sources with a high-frequency response are preferred. Droopers. A constant-current, a drooper, power source displays V-A curves such as those in Fig.(a) (curves A and C). A

considerable negative slope is evident. The internal impedance of the power source is said to be high. Large variations in

the arc (load) voltage (for example, below 40 V) do not appreciably affect the selected current. Droopers are noted for their use in the SMAW process, where the operator sets current to establish a suitable burn-off rate for the selected electrode. Manipulation, that is, shortening or lengthening the arc, does not greatly affect the arc current. However, if welding is conducted on a flatter V-A curve (see curve B in Fig. 3a), then electrode manipulation may change the current considerably. Thus, the operator gains more local control of the current through eye-hand coordination. Increasing the arc length by "pulling back" on the electrode causes the operating point to shift up slightly in terms of voltage, but decrease considerably in terms of current. Open-circuit voltage can be field adjusted in generator types, which provides the operator with a great degree of control through electrode manipulation. Skilled operators use this flexibility to their advantage when performing difficult out-ofposition welds. Constant-Voltage Sources. Equipment designed for constant-voltage processes will display a "stiff," nearly horizontal V-A curve (curve E, F, and G in Fig. 4b). The open-circuit voltage is 43 V for curve F. When serving a 300 A load, it falls to 35 V. This is expressed as: (43 − 35) 8V 2.6V = = 300 A 300 A 100 A

(EQ 4)

Curve H of Fig. 4(b) is representative of an inverter source in the constant-voltage mode. When set to 30 V, the electronic regulation system keeps the output very constant out to 300 A under a static (resistive) load, beyond which it falls abruptly to 0 V at 400 A. Curve K in Fig. 4(b) represents an extended output that may be available on a short-term (dynamic) basis from inverters. The term "constant" is taken to mean relatively constant. Some droop in voltage is permitted, and is even preferred in certain welding situations. Load current at a given voltage is responsive to the rate at which a consumable electrode is fed into the arc. The designer ensures that terminal voltage is maintained during heavy loads by providing tight coupling between the primary and secondary windings of the transformer, or by providing electronic regulation systems, or both. Movable Magnetic Cores. Welding power must be delivered to the arc under complete control of the operator. In the

simplest form, conversion to the lower voltage and selection of current are incorporated into a single, special welding transformer. Movable coils (varying the coupling between primary and secondary windings) and moving iron shunts provide stepless adjustment of the output (Fig. 2a). Saturable Reactors. In heavy industrial equipment, control is imparted by saturable reactors. These reactors comprise

laminated iron core assemblies fitted with power and control windings. A small control current varies the effective magnetic permeability of the iron core and controls flux to the power winding. Considerable gain is obtained. Typically, 5 A of control will command several hundred amperes of welding current. Both ac and dc outputs are available. Direct current output is obtained by the addition of secondary rectifiers. By combining rectifier diodes with self-saturating magnetic cores, a wider range is obtained, referred to as magamp control. Alternating current and ac/dc power sources are manufactured in sizes that range from 180 A Class III "farm welders" to units designed for industrial use and rated in hundreds of amps. Principal attributes are low cost and simplicity. Alternating current is specified wherever arc blow is a concern. It is also preferred for welding aluminum. Direct current machines are favored for most SMAW and GTAW processes, and for essentially all GMAW processes. Silicon-Controlled Rectifiers. In addition to the transformer taps, moving coils, moving shunts, and saturable reactors

mentioned earlier, there are a number of other means by which output can be controlled. Thyristors, commonly known a SCRs, are frequently employed as control elements, replacing magnetic cores. The SCR is similar to a rectifying diode, but is equipped with an additional gating element. With properly phased lowpower gate signals, a configuration of SCRs will deliver hundreds of amperes of current. The output of such a phase-controlled arrangement (Fig. 2b) can be tailored to perform as a stiff (that is, flat slope) constant-voltage source or as a drooping constant-current source through closed-loop feedback controls. When designed for three-phase operation, the dc output ripple frequency is either three times the line frequency (using three SCRs) or six

times, using six SCRs. Output ripple is "smoothed" with an appropriate inductor. Single-phase versions display the ripple of twice the line frequency and require a larger output inductor, L. Because welding current is analog in nature, it has traditionally been adjusted with analog circuitry via rheostats and potentiometers. Recent advances allow digital control with microprocessors and dedicated software; tight-tolerance repeatability is the driving force. Inverters. Inverter welding power sources are smaller and more efficient, and they offer better performance than their

line-frequency counterparts. The size and weight of iron-core components are reduced when operating at higher frequencies (Fig. 5).

FIG. 5 300 A RATED CONVENTIONAL LINE-FREQUENCY POWER SOURCE (RIGHT) AND INVERTER POWER SOURCE WITH SAME RATING (LEFT). COURTESY OF MILLER ELECTRIC MANUFACTURING COMPANY, INC.

With reference to the inverted shown in Fig. 2(c), the incoming utility line directly feeds a rectifier for conversion to direct current, which powers an inverter stage operating in the range from 2 to 100 kHz. Various topologies include forward, buck, series capacitor, and series resonant inverters. Control methods employ either pulse width modulation (PWM) or frequency modulation (FM). A variety of semiconductors are used as switching elements, including transistors (for example, field-effect transistors [FETs] and insulated-gate bipolar transistors [IGBTs]). The transformer, which is often designed of ferrite core material, operates at high frequencies without introducing high losses. The turns ratio determines welding voltage, usually 80 Vrms ac at the secondary winding. A second rectifier converts it back to dc for welding. A small output inductor is required to smooth the high-frequency ripple. Inverters have the greatest application where portability or advanced performance is required. Fast response permits current with fast transitions, low overshoot, and faithful reproduction of pulse programs. Variable Polarity. High-technology power sources permit adjustment of the ac waveform to vary the ratio of EP to EN

current and terms of both width and amplitude (Fig. 3e). Among these are switched dc types, wherein two dc constantcurrent power sources are alternately switched on and off, and arc current is of one polarity and then the other polarity. Almost unlimited frequencies can be generated, with rapid zero crossings. Cycloconverter power sources constitute a family that uses three-phase inputs to generate low-frequency ac for the

GTAW process. Through careful selection of specific portions of each line-frequency cycle, a pseudo square wave is assembled at a frequency lower than line frequency. Rotating Equipment. For a variety of reasons, rotating generators and alternators are often used to supply welding

power. In what is known as a motor-generator (M-G) set, an electric motor (Fig. 2d) drives the shaft of a generator, which also isolates weld power from the input utility power. The flywheel effect of the M-G set ensures stable output.

Where no electric utility power is readily available, engine-driven equipment can supply welding power (Fig. 2e). Gasoline and diesel fuels predominate. Both air- and water-cooled prime movers are available. Some include an auxiliary air compressor for air tool use. Generators produce direct current. A wire-wound armature spins in a frame arranged with electromagnetic field poles. The armature windings cut through these fields cyclically, generating ac. This is collected by commutator, that is, copper segments fitted to the armature shaft and equipped with fixed carbon brushes. The commutator serves as a "mechanical rectifier," converting the armature currents to dc. Both the OCV and current can be user adjusted, providing good control of the V-A slopes. The output is relatively pure dc, with very little ripple content. Engine speed is governor controlled. Most generators provide constant-current output for SMAW and GTAW processes, but adapters are often available for constant-voltage operation in the GMAW process. Auxiliary power for lights, tools, and so on derived from auxiliary stator windings. Alternators (that is, ac generators) produce alternating current, as the name implies (Fig. 2e). A spinning rotor winding is charged with direct current obtained from the exciter and applied through slip rings, creating a rotating magnetic field. No commutator is required. The rotating field cuts through a surrounding stator assembly, the windings of which collectively produce alternating current, either single phase or three phase, as designed. In single-phase designs, the resulting current can be used for ac welding. Output selection is provided in course ranges with reactor taps (Fig. 6), and fine adjustment is provided by rheostatcontrolled current in the revolving field windings. Engine speeds are generally either 1800 or 3600 rev/min, corresponding with four-people and two-pole alternators, respectively. Some alternators eliminate the revolving field winding by employing rotating permanent magnets. If direct current is required, a static (silicon diode) rectifier is included. The combination is referred to as a dc generator-rectifier.

FIG. 6 SCHEMATIC SHOWING TYPICAL ENGINE-DRIVEN WELDING POWER SOURCE

The recent evolution to solid-state controls has raised the performance level of engine-driven machines. Accessories include remote controls, idle devices to reduce engine speed when not welding, battery charging capabilities, and highfrequency arc igniters. Pulsed Power Supplies Pulsed GTAW and GMAW processes require a power source with the ability to deliver precisely shaped current pulses superimposed on a lower background (arc-sustaining) level. In addition to thyristors in phase control, inverters are also used to generate pulses and are especially versatile for such applications. Add-on "pulse formers" have also been designed for attachment for conventional dc welding power sources. Pulse Repetition Rates (PRR). For the GTAW process, PRRs (the number of times per second that a pulse is

transmitted) are relatively low (typically, 1 to 10 PRR). Fast rise times are not essential. However, in the GMAW pulse

process, repetition rates typically range from 100 to 200 PRR. Rise (and fall) times must be short (50 to 500 overshoot must be low to preserve pulse fidelity.

s) and

The mean current of a train of rectangular pulses is given by:

IAV = FTP(IP - IB) + IB

(EQ 5)

where Iav is the average current, IB is the background current, IP is pulse peak current, f is frequency, and TP is pulse width. From Eq. 5, with background current, pulse width, and peak current selected, the average current will be proportional to frequency. Increasing and decreasing pulse frequency in proportion to wire feed speed will balance the current with the burn-off rate of the electrode, maintaining the proper arc length for the GMAW process. Individual droplet size will remain essentially constant. Some pulse power sources generate pulses at line frequency or multiples thereof and stabilize the process by adjusting the power envelope of each pulse. (In a refined development, designated "one-knob control," the operator need only select a wire and gas type; pulse parameters are established by algorithms in the controller. As wire feed speed is adjusted, all other functions "track.") Synergic Pulsed GMAW. A step higher in technology results in synergic pulsed GMAW. Inverters (Fig. 2c) and

secondary switchers (Fig. 2f) ensure process stability by controlling melting rate through modulation of pulse shape and frequency. The relationship between digitally monitored wire feed speed and pulsation parameters is managed by microprocessor look-up tables and similar control strategies. The power source is simply required to exhibit fast response and high peak current capability. Multiple Operator Power Sources On large job sites with a number of arcs, a multiple operator (MO) power source can serve several operators. Diversity of use suggests that all arcs will not be operating simultaneously and that a 0.25 factor is typical. Accordingly, a single, large, 76 V constant potential power source can supply up to 20 SMAW operators. Each operator is equipped with a resistor "grid" to adjust arc current to his particular needs. An alternative MO concept provides individual magamp control modules (typically, eight) fed from a single, large, multiwinding transformer. Each operator can select his preferred polarity and current. Modules can be connected in parallel for greater output. Power Sources for Welding F. James Grist, Miller Electric Manufacturing Company, Inc.; William Farrell and Glen S. Lawrence, Ferranti Sciaky

Resistance Welding Power Sources The function of the power source for resistance welding (RW) is to deliver a predetermined amplitude of current to the welding electrodes that are clamping the workpiece. Current flows from the power source to one electrode, through the workpiece to the opposing electrode, and returns to the power source. The electrodes concentrate current into a small, usually circular area. The flow of this constricted current generates heat, concentrated largely at the interface(s) of the workpiece details to be joined. Heat Input. The rate of heating must be sufficiently intense to cause local electrode/workpiece interface melting. The

opposed electrodes apply a pressure and, when sufficient melting has been achieved, the current is interrupted. Electrode force is maintained while the molten metal solidifies, producing a sound, strong weld. If both the local resistance of the workpiece and the welding current magnitude were constant, then the total quantity of heat, Q, developed in the workpiece would be given (in joules) by:

Q = I2 RT

(EQ 6)

where I is the effective value of current in amperes, R is the resistance of the workpiece in ohms, and t is the duration of flow of current in seconds. However, the resistance is not constant. Contact resistances decrease in magnitude and the bulk resistance of the work column increases as its temperature rises. Furthermore, resistance variations of the workpiece may cause variations in the current amplitude, depending on the nature of the power source. Under these circumstances, the total quantity of heat, Q, is given by: t

Q = ∑ i ² r.dt

(EQ 7)

0

where i is the instantaneous value of current expressed as a function of time and r is the instantaneous resistance that is changing with time. Equipment Selection. The design of resistance welding machines resembles that of a "C" frame press (Fig. 7). The

power source transformer secondary winding is connected to the electrodes on cantilevered arms, one of which provides for adjustable electrode force. Throat depth is the distance from the frame to the electrodes. Throat gap is the vertical distance between the two parallel secondary conductors. The total area of the secondary circuit (throat depth times throat gap) exerts a profound influence on output performance. Any increase in inductance, whether due to geometry or magnetic materials within the secondary, will degrade performance and lower the power factor.

FIG. 7 SCHEMATIC SHOWING PRIMARY COMPONENTS OF A RESISTANCE WELDING MACHINE

The input to resistance welding power sources is usually 230 or 460 V utility power delivering single-phase or threephase alternating current at 60 Hz. All systems require transformers to step down the line voltage to a relatively low value, with a proportional increase in current. The simplest systems consist of a transformer connected to single-phase power. More advanced sources produce a dc welding current through rectification. Rectification Systems. In primary rectification systems, a special three-phase transformer is designed with a very large iron core. Each dc pulse is limited in duration and each is reversed in polarity to prevent saturation of the core. In the majority of applications, a weld is produced by a single impulse of direct current. For heavy gages, pulsation welding is used. A series of dc impulses of alternating polarity is separated by "cool time." The result is a very low frequency alternating current. Hence, the term frequency converter is commonly applied to this type of power source. Figure 8 shows a three-phase half-wave primary rectified power source.

FIG. 8 CIRCUIT DIAGRAM OF THREE-PLATE HALF-WAVE RESISTANCE WELDING POWER SOURCE WITH PRIMARY RECTIFICATION

Secondary rectification systems are available in single-phase full-wave, three-phase half-wave, and three-phase full-wave forms; the latter is most commonly utilized. Three-phase full-wave systems consist of three single-phase transformers, with rectifiers connected to the secondary windings (Fig. 9). With this system, current pulse time is limited only by thermal considerations and polarity is fixed.

FIG. 9 CIRCUIT DIAGRAM OF THREE-PHASE FULL-WAVE RESISTANCE WELDING POWER SOURCE WITH SECONDARY RECTIFICATION

The energy efficiency of a resistance welder depends on the efficiency of delivered energy at the electrodes. Energy is dissipated in the form of heat, because of resistance in the conductors. The current, I, which produces the heat of welding, is simply the secondary voltage divided by the total machine impedance, Z, plus that of the weld: i=

E Z

(EQ 8)

where Z is the vector sum of the resistance and the reactance: Z = R ² + X ² = R ² + (ω L)²

(EQ 9)

where R is the total resistance of the circuit and X = ωL is the total reactance of the circuit (where ωequals 2πf and L is the inductance). Resistance welders designed for single-phase utility power are lowest in investment cost, but highest in energy cost. The high-magnitude alternating current results in large inductive losses and a low power factor. Three-phase primary rectification and three-phase full-wave secondary rectification power sources reduce the power demand to about one half and operate at a higher power factor.

The rapidly pulsing weld zone temperature inherent with a single-phase source is least desirable from a welding standpoint. In addition, copper electrodes suffer from clamping impact, force, and heat, causing the occurrence of "mushrooming," which lowers current density. Mushrooming also reduces resistance in the current path. Welding current will increase, because current is controlled indirectly in resistance welding machines by adjusting secondary voltage (that is, advancing or retarding the firing angle). From Eq 8, a decrease in impedance causes an increase in the current. Therefore, all resistance in the work path influences the current magnitude produced by a given secondary voltage. This influence is small if total system impedance is high, but significant if the system impedance is low. A three-phase primary rectification system exhibits the greatest degree of this self-compensating effect, followed to a somewhat lesser degree by the three-phase secondary rectification category. Finally, a thermoelectric effect is associated with passing current from a copper alloy electrode to a workpiece of a given material. This phenomenon is important when welding aluminum or magnesium alloys with a single impulse. Reversing the current polarity for each successive weld, an inherent characteristic of the three-phase primary rectification system, provides the maximum electrode life between dressings. Duty Cycle. Resistance welding is inherently an intermittent process with a series of very short current periods followed

by "off" or "cool" periods. Therefore, this type of power source is rated on a 50% duty cycle and a 1 min integrating period. The relationship between kVA rating at 50% duty cycle and kVA demand at actual duty cycle is as follows: KVA50% = KVADemand 2.( D.C.)

(EQ 10)

where D.C. is the duty cycle in %, entered in the form of a decimal. Duty cycle is defined by: D.C. =

N .T 60. f

(EQ 11)

where N is the number of welds per minute, T is the weld time in cycles, and f is the frequency in hertz. Heat Controls. Resistance welds are short and are measured in cycles. Because timing is very important, digital

microprocessor controls abound. Electrochemical contactors have been replaced by SCRs or ignitrons connected in inverse parallel to handle the high currents. Coarse heat adjustment is accomplished by transformer tap switch setting, and fine adjustment is achieved through phase shifting, with digital circuitry providing the highest repeat accuracy. Additional information on resistance welding is available in the Section "Joining Processes" in this Volume. Power Sources for Welding F. James Grist, Miller Electric Manufacturing Company, Inc.; William Farrell and Glen S. Lawrence, Ferranti Sciaky

Electron-Beam Welding Power Sources Electron-beam welding (EBW) requires a variety of power supplies in two categories: those that generate and control beam position and those that control beam power. The outputs of various power supplies are coordinated by a machine control system. Beam Position and Intensity. Beam position arrangements for alignment, deflection, and focus are dominated by magnetic devices in electron-beam welders. Because the magnetic field effect on the beam (the magnetomotive force) is directly proportional to the current in the windings of the device, these power supplies are typically current regulated.

MMF = N · I

(EQ 12)

where N is the number of turns in the coil and I is the current in the coil. The desired current value is programmed by potentiometer, computer, or function generator.

A second factor regarding the effectiveness of magnetic devices is the velocity of the electrons in the beam, as determined by the accelerating voltage. At the voltages used in EBW (typically, 50 to 150 kV), relativistic effects must be considered. In situations where the beam potential is changing, but it is desired to have the effect of the magnetic devices be constant, then the effective magnetic excitation, k, must be constant: K² =

( N .I )² V*

(EQ 13)

where NI is the magnetomotive force and V* is the relativistic voltage. The relativistic voltage can be expressed in volts as:

V* = V + (0.9785 × 10-6)V2

(EQ 14)

Sensors monitor the beam potential generated by the high-voltage (HV) supply and correct the control signal to the magnetic device power supply. Thus, the action of the device can be held constant as beam potential changes. This correction can be achieved by using computer software or an analog circuit. Such correction is often called square root compensation, because of the relationship indicated in Eq 13. Alignment coils are used for slight corrections to the x and y positions of the beam as it enters the magnetic focusing coil. Power supplies for these devices are typically potentiometer programmed by the operator in an alignment procedure. These supplies are characterized by a high degree of long-term stability, with minimal need for dynamic response. Selected installations employ alternate means for achieving alignment. The magnetic focus coil acts very much as a glass focusing lens would for optics. It takes a divergent electron beam and converges it at a more distant point along the beam axis. The beam focus point from the magnetic lens is roughly inversely proportional to the square of the lens excitation: f ∝

V* ( N .I )²

(EQ 15)

The magnetic lens supply is usually controlled by computer output in order to coordinate beam focus with the weld program. Manual control (operator offset) and square root compensation are also usually provided. These supplies are of high stability and moderate dynamic response. The deflection coil moves the beam along the x and y axes to position the beam over the surface of the workpiece. Deflection rates can range from dc to high audio frequencies. Thus, the deflection supply is a high-quality currentregulated stereo amplifier. Amplifier input controls allow computer input for coordinated movements of the part and beam, function generator inputs for beam deflection patterns (rasters, and so on), and dc offsets to patterns. These are highly stable current-regulated coils drivers with wide dynamic response. High-Voltage Control. The electron beam is generated in an electron gun (resembling a diode or triode of cylindrical symmetry) that generates a paraxial flow of electrons. The power in an electron beam comes from the high-voltage supply that generates the large voltage difference between the anode and cathode of the gun. Electrons are emitted from the cathode and, accelerated by the voltage, fall to the anode. The beam passes through a central aperture in the anode and drifts through various alignment, focus, and deflection elements to strike the workpiece (Fig. 10). Both the workpiece and the anode are typically at ground potential, whereas the cathode is highly negative.

FIG. 10 SIX INDIVIDUAL POWER SUPPLIES THAT CONTROL POSITION AND INTENSITY OF THE BEAM IN EBW PROCESS

Control and regulation of the high voltage is usually done on the input to the high-voltage transformer/rectifier oil-filled tank. The input control installations range from motor-driven variable autotransformers, motor-alternators with field controls, SCR switching or phase control, to power MOSFET (metal-oxide semiconductor field-effect transistor) and IGBT devices in various switch mode regulators. In a three-phase input unit, a standard technique is to place the switching control in the primary neutral conductor of the HV transformer (Fig. 11). The secondary circuitry is three-phase or six-phase, full-wave rectified, and filtered.

FIG. 11 CIRCUIT DIAGRAM OF A HIGH-VOLTAGE CONTROL IN NEUTRAL IN A THREE-PHASE INPUT UNIT

In an alternate arrangement, line voltage is converted to dc, which powers an inverter to produce high-frequency singlephase input to the high-voltage transformer. Control can be provided either by regulating the level of the dc supply or by controlling the inverter pulse width or frequency. The HV transformer can be a conventional step-up unit or a voltage multiplier. These supplies can be characterized as voltage regulated and input controlled, with a moderate response rate. Thermionic Emission. The cathode heating unit provides the energy to raise the cathode to thermionic electron emission

temperatures. This is usually a dc source. For beam-heated cathodes, discharge-heated cathodes, and others, it may be more complicated, to a point of supplying power to a second miniature gun. In a space-charge limited electron beam, the cathode temperature must be held high enough that temperature-limited emission (Eq 16) is greater than space-charge limited flow (Eq 17):

IP = AT2 E −eφ / kt IP = K(VP + µ VG)3/2

(EQ 16) (EQ 17)

In Eq 16 and 17, Ip is beam current, Vp is accelerating voltage, A is Dushman's constant, Vg is bias voltage, and T is temperature. The bias supply provides voltage to the control electrode and regulates the magnitude of the beam current. Bias supplies can be input controlled with ground-level sensing of beam current on the return of the HV supply, or the beam can be sensed and regulated at high voltage using optical isolation (Fig. 12).

FIG. 12 CURRENT FEEDBACK AND BIAS CONTROL AT HIGH-VOLTAGE POTENTIAL

Ground-level sensing systems tend to use high-frequency inputs from linear amplifiers or switch mode inverters in order to achieve response on the order of a millisecond. Submillisecond response bias supplies are achieved by sensing and controlling the beam current at the high-voltage output. Because bias voltage is in the low kV range, vacuum tubes are used for output control. Additional information on electron-beam welding is available in the article "Electron-Beam Welding" in this Volume. Fundamentals of Weld Solidification Harvey D. Solomon, General Electric Company

Introduction OF ALL PHASE TRANSFORMATIONS, few have been more widely observed and studied than the transformation of a liquid to a solid (that is, solidification). The process of solidification is the same in all cases, whether it is the freezing of water on a windshield or in a freezer, or the solidification of metal in a casting or in the weld that joins two solids. The process is controlled by the free energy of the liquid phase, Gl, relative to that of the solid, Gs. This is depicted in Fig. 1, which shows the behavior of a pure (single component) material. Above the freezing temperature, Tf, the liquid phase has the lower free energy and is therefore stable, but below Tf, the solid is the stable phase. At Tf, both phases are in equilibrium, that is, Gl = Gs.

FIG. 1 TEMPERATURE DEPENDENCE OF BULK FREE ENERGY OF THE LIQUID AND SOLID PHASES IN SINGLECOMPONENT SYSTEM. THE SOLID LINE PORTION AND THE DASHED LINE PORTION OF EACH GL, AND GS CURVE INDICATE THE STABLE AND UNSTABLE PHASES, RESPECTIVELY, OF THE FREE ENERGY ON EITHER SIDE OF TF. SOURCE: REF 1

In the transition from one phase to another, the change in free energy, ∆G, is the difference in free energy of the product and the reactant. This free energy change can be expressed in terms of the enthalpy and the entropy changes, that is, for the transformation of a liquid to a solid during freezing:

∆G = GS - GL = (HS - HL) - T(SS - SL) = ∆H – T∆S

(EQ 1)

At the freezing temperature, Tf, ∆G = Gs - Gl = 0, because the free energy of the two phases is the same, and ∆H = Tf∆S, It is necessary to cool below Tf for solidification, because at Tf both the solid and liquid phases are present and in equilibrium. Below Tf, ∆G is not equal to zero (Fig. 1 shows that Gs < Gl) and is given by Eq 1 with T = T', where T' < Tf. Because the ∆H and ∆S are not strong functions of temperatures, they can be assumed to be temperature independent. Therefore, at any temperature, ∆H = ∆Hf and ∆S = ∆Sf, where ∆Hf and ∆Sf are the values of the enthalpy and the entropy changes for the equilibrium reactions at Tf (that is, the latent heat of fusion and the entropy change on fusion, respectively). Combining these enthalpy and entropy expressions, the fact that ∆H = Tf ∆S, and Eq 1, then at T', one obtains:

∆G = (∆HF/TF)(TF - T')

(EQ 2)

where ∆Hf, the latent heat of fusion, is negative. Hence, in agreement with Fig. 1, ∆G is negative. The greater the amount of undercooling (supercooling) below Tf (Tf - T'), then the greater the thermodynamic driving force for solidification. However, even when the conditions of Eq 2 are met, the liquid does not spontaneously transform to the solid below Tf. Rather, small amounts of solid nucleate and grow to produce complete solidification. Nucleation creates a new surface, that is, the surface between the solid and the liquid. The energy per unit are of this surface is the surface tension, γ, which is always positive. For solidification to occur, the increase in energy associated with the surface energy must be balanced by a greater decrease in the free energy of the solid relative to that of the liquid. This requires undercooling, as shown by Eq 2.

There are three ways in which a solid can form: homogeneous nucleation, heterogeneous nucleation, and epitaxial growth. Homogeneous nucleation occurs when there is no foreign body (mold wall, solid particle in the melt, etc.) on which to form the solid. Figure 2 shows the balance of the surface tension and the bulk free energy per unit volume, ∆Gv, as a function of the size of the nucleus that forms during homogeneous nucleation. The ∆Gv, value is just ∆G (as given by Eq 2) divided by the molar volume of the solid, Vs.

FIG. 2 FREE ENERGY OF FORMATION OF A NUCLEUS AS FUNCTION OF ITS RADIUS. SOURCE: REF 1

For a spherical nucleus of radius greater than rc, the volume free energy decrease outweighs the increase in energy that is due to the surface energy, and the nucleus is stable. At r = r*, the net energy is a maximum, but if additional atoms are added to the nucleus, then the energy decreases. The value r* is thus the radius at which the spherical nucleus is just stable, because as additional atoms are added, the energy is decreased. The value r* is given by:

R* = -2γ/∆GV

(EQ 3)

which is a positive number, because ∆Gv is negative. Substituting into Eq 2, one gets:

R* = -2 γTFVS/∆H F(TF - T')

(EQ 4)

The r* is positive because ∆Hf is negative and, for undercooling, (Tf - T') is positive. This equation is a general expression relating the radius of curvature of a surface and the degree of undercooling that is required for solidification of that surface. The greater the degree of undercooling, the smaller the radius of curvature that is stable. This equation will be considered again when nonplanar solidification is discussed. Nucleation requires a thermal energy either equal to or greater than ∆G* in order for a spherical nucleus of radius r* to form. The ∆G* value is the activation energy, which, as Fig. 2 shows, is a positive quantity. The greater the degree of undercooling, the more negative the free energy, the less positive the ∆G* value, and the greater the rate of nucleation of the solid. For homogeneous nucleation, ∆G* is given by:

∆G* = (16/3)(π γ 3TF2VS2)/(∆HF2∆T2) where ∆T = Tf - T'.

(EQ 5)

Heterogeneous nucleation develops when the solid forms on a foreign body. Now, the interaction of the nucleus and the foreign body must be considered. This interaction is defined in terms of the wetting angle between the nucleus and foreign body. The activation energy for heterogeneous nucleation is given by:

∆G* =

2 2 4π γ ³T f Vs (2 − 3cos θ + cos ³θ ) 3 ∆H 2f ∆T ²

(EQ 6)

The angle is 180° if there is no wetting, and Eq 6 reduces to Eq 5, that is, with no wetting, there can only be homogeneous nucleation. With a wetting angle of 90°, ∆G* is one half the value given by Eq 5. Instead of requiring a spherical nucleus of radius r*, a hemispherical nucleus of the same radius is necessary, which requires one half the number of atoms and one half the activation energy. This is the sort of nucleation that develops in a casting, where the mold wall acts as the foreign body. When the wetting angle is zero, Eq 6 shows that the activation energy is zero. This is the case for epitaxial growth on a substrate, where epitaxial derives from the Greek epi, upon, and taxis, to arrange, that is, to arrange upon. In effect, no new surface is being formed. Atoms are just being added to the substrate, thereby extending it. The important point is that no activation energy nor undercooling is required to add atoms onto an existing substrate. This is the situation that develops when a liquid solidifies on a substrate of the same material or one that is similar in composition and structure, as in the solidification of a weld.

Reference

1. J.D VERHOEVEN, FUNDAMENTALS OF PHYSICAL METALLURGY, JOHN WILEY & SONS, 1975

Fundamentals of Weld Solidification Harvey D. Solomon, General Electric Company

Comparison of Casting and Welding Solidification A weld can be thought of as a miniature casting. The fundamentals of weld solidification are the same as those of a casting, but with different boundary conditions. These differences and their effects are described below. First, by its very nature, a sound weld must attach to the metals being joined, whereas a casting must not adhere to the mold wall. To achieve this, the mold is treated to prevent adherence, whereas the joint to be welded is prepared to promote adherence. The practical result of this preparation is that there is generally excellent heat transfer through a weld joint, but relatively poor heat transfer through the mold, because oxides are often used as mold materials or for mold coatings (the molten metal will not wet the oxide and therefore cannot adhere to it). Second, heat is continually being added to the weld pool as it travels, whereas no heat is added to a casting after the pour, except for possibly modest heating of the mold. The practical result of this and of the first consideration described above is that the temperature of the casting is relatively uniform. In contrast, a very large temperature gradient develops in a weld. The center of a weld pool reaches a very high temperature (typically, 2000 to 2500 °C, or 3630 to 4530 °F), which is limited by the vaporization of the weld metal. At the sides of the weld pool, where solidification is taking place, the temperature is the solidification temperature, which is typically about 1000 °C (1800 °F) less than the temperature at the center on the weld pool. Thus, a positive temperature gradient is developed in a weld (that is, measured from the fusion line of the weld, the temperature increases going into the molten material). A casting will typically supercool to below the solidification temperature. Solidification will first develop by heterogeneous nucleation at the mold wall. The latent heat and the poor heat transfer through the mold wall will then cause a temperature rise at the mold wall. Thus, a small negative temperature gradient can develop briefly in a casting. The nature of this gradient is important, because it

promotes nonplanar solidification (to be discussed below) in the casting, whereas the positive gradient that develops in a weld helps to limit nonplanar solidification. Third, solidification is developed at the mold wall of a casting by heterogeneous nucleation, whereas epitaxial growth develops at the weld fusion line. Thus, some supercooling is required for the casting, but only a vanishingly small degree of supercooling is required for the weld. (As was discussed, epitaxial growth requires no supercooling, but some slight degree of undercooling is required to shift the reaction from an equilibrium between the solid and liquid, to all solid.) This undercooling in a weld would typically be less than 1 °C (1.8 °F). In contrast, in castings when there is a lack of effective nuclei for hetergeneous nucleation, undercoolings of 100 °C (180 °F) or more are possible. In general, however, the undercooling in castings is typically only on the order of a few degrees centigrade. Fourth, the generally larger volume of a casting, relative to that of a weld, and the poorer heat transfer makes the cooling rate and the solidification rate much lower for castings than for welds. Fifth, as a casting solidifies, the volume of the remaining liquid decreases. Thus, the shape of the molten pool is continually changing. In a weld, the weld pool shape is generally kept constant as it travels (if the heat input and section geometry are constant). Sixth, because of the stirring action of the arc and the action of Marongoni surface tension gradient induced convection forces, there is good mixing of the molten weld pool. In contrast, there is comparatively little mixing of the molten material of a casting. As discussed in the section "Solidification of Alloy Welds (Constitutional Supercooling)" in this article, all of these factors influence the nature of weld solidification, relative to what is observed in a casting. Fundamentals of Weld Solidification Harvey D. Solomon, General Electric Company

Solidification of Alloy Welds (Constitutional Supercooling) The solidification of an alloy is much more complex than that of a pure metal, whether in a weld pool or in a casting. Figure 3 shows a simple binary isomorphous phase diagram. The temperatures of the liquidus and solidus lines of Fig. 3 can be approximated by straight lines (for a limited temperature range). (The liquidus defines the temperature above which only liquid is present and the solidus the temperature below which only solid is present. The liquidus and solidus lines bound the two-phase liquid plus solid region.) The liquidus line is given by:

TL = TMA + MLCL

(EQ 7)

where TmA is the melting point of pure A, ML is the slope of the liquidus, and CL is the composition of the liquid at TL. The solidus is described by a similar expression.

FIG. 3 BINARY ISOMORPHOUS PHASE DIAGRAM SHOWING THE LIQUID AND SOLID COMPOSITIONS AT SELECTED TEMPERATURES

Another important quantity that can be estimated from the phase diagram is the distribution coefficient, K, which is defined as:

K = CS/CL

(EQ 8)

where CS and CL are the solidus and liquidus compositions, respectively, defined by equilibrium tie lines. Assuming that the solidus and liquidus lines are straight lines means that K is not a function of temperature. In general, the liquidus and solidus lines are curved, so that in actuality, K is a function of temperature. For simplicity, this curvature will be neglected and, instead, limited temperature ranges where K can correctly be assumed to be a constant will be considered. The use of a distribution coefficient defined by equilibrium tie lines implies that equilibrium exists between the liquid and the solid at the solid-liquid interface. This equilibrium need not exist elsewhere. Figure 3 and the following discussion are based on a K value of less than 1 and a negative ML value. The following discussion is also true for diagrams where K is greater than 1 and ML is positive. The distribution coefficient defines how the solute is distributed or partitioned between the solidifying solid and the remaining liquid. The slope of the liquidus line and the distribution coefficient are important in defining the composition gradient that develops when alloys solidify. In general, the distribution coefficient, K, is a function of temperature, but as already noted, to simplify the analysis, it will be assumed that over a limited temperature range, it is a constant. An equilibrium phase diagram, such as that of Fig. 3, assumes that equilibrium is maintained in all of the solid and liquid and that there are no composition gradients in the solid or liquid. It defines the equilibrium solid and liquid compositions at any temperature. Figure 3 shows that the solidification of an alloy of composition C0 will begin when the temperature is reduced below T1. If equilibrium is maintained, then solidification will be complete when the temperature is reduced below T3. A two-phase mixture of solid plus liquid will be present between T1 and T3, with the relative amount of each phase given by the lever rule:  C − Cs 2  %liquid =  0 100%  C1.2 − Cs 2 

(EQ 9)

at T2. The free energy curves corresponding to this equilibrium are much more complex than those of Fig. 1, which is relevant for a pure material.

The equilibrium predicted by Fig. 3 is never realized in practice, because it would require an infinitely slow cooling rate. When considering castings, it is generally assumed that the equilibrium predicted by the phase diagram is maintained at the solid-liquid interface, but that no composition change occurs in the solid as the temperature is lowered. At T1, a solid of composition CS1 forms. CS1 is equal to CL1K, and because at the liquidus CL1 = C0 the CS1 = C0K. When the temperature drops to T2, the solid that forms has a composition of CS2, and so on. A composition gradient is thus developed in the solid. As the solid forms, solute is rejected into the liquid, and it is assumed that this solute is mixed into the liquid, raising its composition to that predicted by the phase diagram. This occurs even though the mixing action in a casting is relatively poor, because the cooling rate is generally low enough to provide enough time to allow for the required solute redistribution by diffusion in the liquid. As a result, no gradient is developed in the liquid. The composition of the solid at the solid-liquid interface is still given by the equilibrium phase diagram. However, the previously formed solid is deficient in solute. As a consequence, complete solidification does not occur at T3. Further cooling is required, with solidification being completed only when the average solid composition is C0. The boundary conditions for a weld are somewhat different from those of a casting. The composition gradient that develops in the solid is assumed to be the same as that which forms in a casting, but the behavior of the liquid is assumed to be different. The mixing of the melt is greater than that developed in a casting, but the cooling rate is much higher. Thus, the solute that is rejected as the solid forms is not completely mixed into the liquid. Rather, the bulk liquid remains at the initial composition, C0, with a solute gradient formed in the liquid near the solid-liquid interface. At the liquids, T1, the solid that forms has a composition of CS1 = C0K. At T2, the solid composition is CS2. As with castings, it is assumed that the solid-liquid equilibrium predicted by the equilibrium phase diagram is maintained at the solid-liquid interface. Therefore, the liquid in equilibrium with this solid is CL2 = CS2/K. However, the lack of complete mixing means that the bulk liquid remains at C0, and that a composition gradient forms directly ahead of the solid-liquid interface. This composition gradient is illustrated in Fig. 4, which shows the composition gradient developed by unidirectional solidification and the temperature gradient driving this solidification. The solidification is modeled by assuming that a bar is being unidirectionally solidified, but this approach is applicable to solidification in general.

FIG. 4 INITIAL STAGE OF DIRECTIONAL SOLIDIFICATION OF A BAR, SHOWING THE TEMPERATURE PROFILE DRIVING THE SOLIDIFICATION AND THE COMPOSITION PROFILE ESTABLISHED FOR AN ALLOY WITH A SOLUTE CONTENT OF C0. SOURCE: REF 4

At temperature T3, the solid that forms has a composition of C0. The solid that is forming now has the same composition as the bulk liquid. Thus, an equilibrium situation exists. Solidification can now continue without having to further reduce the temperature. For castings, however, the composition gradient in the solid and complete mixing of the solute into the

liquid leaves the liquid still present at T3, thereby reducing the liquidus temperature. This means that, in a casting, the temperature must be lowered even more before further solidification can take place. Figure 4 shows the situation during weld solidification, whereas Fig. 5 shows the composition when solidification is complete. For most of the solidification, the solid forms at C0. The exceptions are at the start and at the end of solidification. The first solid that forms does so with a solute composition of CS1, which the phase diagram shows is less than C0. A composition gradient develops until temperature is lowered to T3, when the solid composition reaches C0 and an equilibrium is reached between the solid being formed and the composition of the bulk liquid. At the end of solidification, there is an increase in the solute content. This occurs to accommodate the solute buildup at the solid-liquid interface. At the end of the solidification, there is no longer any bulk liquid at C0. The solute that is ahead of the solidliquid interface raises the composition of the remaining liquid as it is rejected into this final small volume of liquid. This obviates the assumed boundary condition (a lack of mixing of the solute into the bulk liquid), and the situation becomes more like that of a casting (complete mixing in this small liquid volume). The final liquid composition of this final transient is the terminal composition predicted by the phase diagram, that is, the final liquid is driven to the composition of the lowest melting point. In the case of the simple binary isomorphous diagram used for this example, this is pure B. For a simple binary eutectic, it would be the eutectic composition.

FIG. 5 FINAL COMPOSITION PROFILES FOR THE UNIDIRECTIONAL SOLIDIFIED BAR SHOWN IN FIG. 4. THE INITIAL AND FINAL TRANSIENT ARE SHOWN FOR K < 1. R, RATE OF SOLIDIFICATION; DL, DIFFUSION RATE OF THE SOLUTE IN THE LIQUID. SOURCE: REF 4

The profile of the liquid composition gradient that is ahead of the solid-liquid interface is given by:

CL = CA e − [ RX L / DL ] + C0

(EQ 10)

where R is the rate of solidification. DL, is the diffusion rate of the solute in the liquid, Ca is the increase in the composition of the liquid at the solid-liquid interface (relative to C0), and XL is the distance into the liquid that is ahead of the solid-liquid interface. The liquid at the solid-liquid interface has a composition of CL3, which is in equilibrium with CS3 = C0. Therefore, CL3 = C0/K. The enrichment of the solute in the liquid, Ca, is thus (C0/K) - C0 or C0(1 - K)/K. The distribution coefficient, K, is thus an important variable in determining the composition profile in the liquid and, as will be shown, in determining the nature of the solidified microstructure. Substituting Eq 10 into Eq 7, one gets the expression describing the temperature variation of the liquidus that is ahead of the solid-liquid interface, namely:

TL = TMA + ML CA e − [ RX L / DL ] + ML C0

(EQ 11)

Figure 6 illustrates this liquidus profile, along with three possible temperature gradients that could actually exist in the liquid. They are all positive, reflecting the positive gradient that develops in a weld. Gradient A is greater than the slope of the liquidus at the solid-liquid interface, but those of B and C are not. The slope of the liquidus can be found by differentiating Eq 11 at XL = 0, that is, at the solid-liquid interface. For the temperature gradient to be less than this slope, the gradient G must be:

G < MLRC0(K - 1)/DLK

(EQ 12)

FIG. 6 LIQUIDUS PROFILE OF A SOLID-LIQUID INTERFACE. (A) LIQUIDUS PROFILE FOR STEADY-STATE CONSTITUTIONAL SUPERCOOLING. (B) THREE LIQUIDUS TEMPERATURE GRADIENTS (CURVES A, B, AND C) COMPARED TO THE LIQUIDUS PROFILE. SOURCE: REF 6

Equation 12 is important because it describes the temperature gradient required for constitutional supercooling (Ref 3, 4, 5). The establishment of the composition gradient in the liquid that is ahead of the solid-liquid interface produces a variation in the temperature of the liquidus. For curves B and C, this has allowed the actual temperature to be below the liquidus (that is, supercooling has developed, even though the temperature of the bulk liquid is above the bulk liquidus temperature). This is constitutional supercooling, and as will be shown, this promotes nonplanar solidification, even though there is a positive temperature gradient. The liquidus profile shown in Fig. 6 is correlated with the propensity for hot cracking. The greater the temperature difference between that of the solid-liquid interface and that of the bulk liquid, the greater the tendency for hot cracking. This is discussed in more detail in the article "Cracking Phenomena Associated with Welding" in this Volume. The constitutional supercooling model for describing weld solidification is presented because it qualitatively describes the evolution of different microstructures (described in the section "Development of Weld Microstructures" in this article). Many newer nonequilibrium approaches to solidification give better quantitative results. These approaches will be

discussed in the section "Nonequilibrium Effects: High-Rate Weld Solidification and Composition Banding" in this article.


3. J.A. BURTON, R.C. PRIM, AND J. SLICHTER, J. CHEM. PHYS., VOL 21, 1953, P 1987-1991 4. W.A. TILLER, K.A. JACKSON, J.W. RUTTER, AND B. CHALMERS, ACTA METALL., VOL 1, 1953, P 428-437 5. W.A. TILLER AND J.W. RUTTER, CAN. J. PHYS., VOL 34, 1956, P 96-121 6. W.F. SAVAGE, E.F. NIPPES, AND T.W. MILLER, WELD. J., VOL 55, 1976, P 165S-173S Fundamentals of Weld Solidification Harvey D. Solomon, General Electric Company

Development of Weld Microstructures Nonplanar solidification develops when a protrusion moves ahead of the rest of the solid-liquid interface and continues to grow in a stable manner. This increases the surface area. The stable radius is given by Eq 4, which shows that the protrusion must move into a supercooled region in order to be stable. This is possible in a casting, because a negative temperature gradient is developed (that is, the temperature decreases from the solid-liquid interface into the liquid). In a pure material (single-component system), the positive temperature gradient that is established in the liquid of a weld prevents this nonplanar growth. In an alloy, however, constitutional supercooling allows nonplanar solidification in a weld, even with a positive temperature gradient. The lower the gradient, G, (compare curve C with curve B in Fig. 6), the greater the degree of constitutional supercooling. This is illustrated in Fig. 7, which shows several temperature gradients (Fig. 7f) compared to the variation in the liquidus and the nature of the solid-liquid interface (Ref 5, 6, 7). For a curve a, the temperature gradient is steeper than the liquidus temperature curve and only planar solidification is possible. For curve b, there is a shallower gradient and the solid-liquid interface can move by cellular growth. The gradient is successively lower for c and d, which allows larger protrusions and the development of cellular dendritic and columnar dendritic growth. At curve e, the gradient is so shallow and the resulting liquidus temperature is so far above the actual temperature in the liquid that equiaxed dendrites can form ahead of the solid-liquid interface.

FIG. 7 SCHEMATICS SHOWING MICROSTRUCTURE OF SOLID-LIQUID INTERFACE FOR DIFFERENT MODES OF SOLIDIFICATION AND THE TEMPERATURE GRADIENTS THAT GENERATE EACH OF DIFFERENT MODES. (A) PLANAR GROWTH. (B) CELLULAR GROWTH. (C) CELLULAR DENDRITIC GROWTH. (D) COLUMNAR DENDRITIC GROWTH. (E) EQUIAXED DENDRITE. (F) FIVE TEMPERATURE GRADIENTS VERSUS CONSTITUTIONAL SUPERCOOLING. SOURCE: REF 6

Figure 8, which is based on experimental observations, relates the gradient, solidification rate, and composition to the type of structure developed (Ref 5, 6, 7). The basis for this figure lies in Eq 12. If R is moved to the left side of Eq 12, then the right side contains only material variables. For a given alloy system (where DL, ML, and K are fixed), the greater the C0, the greater the tendency for constitutional super-cooling and nonplanar solidification. Equation 12 predicts that the

smaller the G/R, the more nonplanar the solidification. Figure 8 shows that, experimentally, the correlation is with G/ R , rather than G/R.

FIG. 8 PLOT OF SOLUTE CONTENT VERSUS SOLIDIFICATION PARAMETER TO SHOW THE NATURE OF SOLIDIFICATION. SOURCE: REF 6

As the dendrite, or columnar grain, grows into the liquid, it should follow the behavior shown in Fig. 5. (As will be discussed in the section "Nonequilibrium Effects: High-Rate Weld Solidification and Composition Banding" in this article, the experimentally determined compositions do not always exactly agree with that predicted by Fig. 5.) The center of each cell or dendrite behaves like the initial transient, that is, the solute content is low. The regions between the cells or dendrites behave like the final transients, that is, they are enriched in solute. This is illustrated schematically in Fig. 9. Cellular growth occurs for a smaller C0 and/or for a larger G/ R than that for dendritic growth (see Fig. 8).

FIG. 9 SCHEMATIC SHOWING SOLUTE DISTRIBUTION AT THE DENDRITE OR CELL CORE AND IN THE INTERCELLULAR OR INTERDENDRITIC REGIONS. (A) CELLULAR GROWTH. (B) DENDRITIC GROWTH. SOURCE: REF 8

The constitutional cooling described by Fig. 4, 5, 6, 7, 8, and 9 occurs on a very fine scale. Solving Eq 10 shows that the solute composition that is ahead of the solid-liquid interface will decrease to almost C0 when XL = 5DL/R. For a typical liquid diffusion rate of 5 × 10-5 cm2/s (5 × 10-8 ft2/s) and a solidification rate of 5 mm/s (0.2 in./s), the liquid composition drops to C0 within about 5 μm (200 μin.). In the initial transient, the distance over which the solid composition increases to C0 is given by 5DL/RK. Thus, for K < 1, this distance is greater than the region in the liquid just ahead of the solidliquid interface. This is illustrated in Fig. 5. This difference is also important in defining the extent of the final transient, which is of the order of the length of the zone in the liquid where the composition drops to C0. The length of the initial transient is thus 1/K larger than that of the final transient. The length of the solute-poor cell or dendrite core is thus 1/K larger than that of the solute-rich intercellular or interdendritic regions. The solute-rich region between the cells or dendrites has a lower freezing temperature than that of the cellular or dendritic core. Thus, it is possible to have liquid films between impinging cells or dendrites. These films cannot support the thermally induced shrinkage stress that develops during solidification, and hot cracking can develop (see the article "Cracking Phenomena Associated with Welding" in this Volume). The hot cracking produced by nonplanar solidification can be minimized by reducing C0 (that is, by using purer materials) and/or by increasing G/ R (see Fig. 8). The temperature gradient in the liquid, G, is inversely proportional to the heat input, Q, which is given by:

Q = FEI/V

(EQ 13)

where f is the efficiency of the welding process, E is the arc voltage, I is the arc current, and V is the welding torch velocity. Thus, G can be increased by decreasing either the arc voltage or arc current, or by increasing the torch velocity. The influence of the velocity is counterbalanced by the fact that the solidification rate, R, is proportional to V (see the section "Effect of Welding Rate on Weld Pool Shape and Microstructure" in this article). Thus, increasing V increases R, which tends to increase the degree of constitutional supercooling.


5. W.A. TILLER AND J.W. RUTTER, CAN. J. PHYS., VOL 34, 1956, P 96-121 6. W.F. SAVAGE, E.F. NIPPES, AND T.W. MILLER, WELD. J., VOL 55, 1976, P 165S-173S 7. F. MATSUDA, T. HASHIMOTO, AND T. SENDA, TRANS. NATL. RES. INST. MET. (JPN), VOL 11, 1969, P 43 8. W.F. SAVAGE, C.D. LUNDIN, AND A.H. ARONSON, WELD. J., VOL 44, 1965, P 420S-425S Fundamentals of Weld Solidification Harvey D. Solomon, General Electric Company

Effect of Welding Rate on Weld Pool Shape and Microstructure The velocity of the welding torch affects not only the rate of solidification, but the shape of the weld pool and the propensity to develop centerline hot cracks. The shape of the weld pool is dictated by the velocity, V, at which the welding torch moves and by the rate at which heat can be removed at the solid-liquid interface. To keep a constant shape, the rate of new melting must be exactly balanced by the solidification rate. On the average, this solidification occurs normal to the solid-liquid interface, because this is the direction of the maximum temperature gradient and, thus, the direction of the maximum heat removal. Simple geometry requires that for the weld pool shape to remain constant, the solidification rate, R, must be related to the torch velocity by:

R = V COS φ

(EQ 14)

where φ is the angle between the normal to the solid-liquid interface and the direction of motion of the torch (Fig. 10). At the back of the weld pool, where φ = 0, R = V. If this were not so, then the weld pool would be elongating or shortening. At the side of the weld pool, where φ = 90°, R = 0. If this were not the case, then the weld pool would be decreasing in diameter. With a constant weld pool shape, solidification at the side of the weld pool occurs as the region adjacent to where φ = 90° sweeps by, rather than by the motion normal to the weld pool velocity.

FIG. 10 SCHEMATIC OF A MOVING WELD POOL SHOWING THE RELATIONSHIP BETWEEN VELOCITY OF TRAVEL OF WELDING TORCH, V, AND THE RATE OF SOLIDIFICATION, R, AT SELECTED POINTS ALONG WELD POOL BOUNDARY. SOURCE: REF 9

Equation 14 predicts that R will vary around the weld pool. Figure 8 shows that as R varies, so does the microstructure. The result of these two effects is a variation of the microstructure around the weld pool (Ref 7). Planar solidification is to be expected at the sides of the weld pool, with the extent of the nonplanar solidification increasing towards the center of the weld pool. This is shown schematically in Fig. 11.

FIG. 11 SCHEMATIC SHOWING VARIATION OF MICROSTRUCTURES IN RESPONSE TO VARIATION OF THE SOLIDIFICATION RATE AROUND THE WELD POOL. SOURCE: REF 7

The coarseness of the nonplanar structure is also a function of R. The dendrite tip radius varies inversely with the square root of R. Not only does the dendrite tip radius become smaller as R increases, the dendrite arm spacing does also, although the exact form of the dependency is more complex. The same general inverse dependency with R is also true for the cell spacing when the solidification is cellular. The ability of the solid-liquid interface to move at a velocity, V, depends on two factors: the maximum solidification rate, which is a function of the crystallographic orientation of the solid at the point in question, and the ability of the temperature gradient in the solid to remove the heat and establish the appropriate temperature for solidification to proceed. Consider first the influence of the orientation of the solid grain at the solid-liquid interface. For body-centered

cubic (bcc) and face-centered cubic (fcc) metals, the preferred growth direction for epitaxial solidification is the direction. For hexagonal close-packed (hcp) metals, the preferred growth direction is . The planar growth of grains oriented with this direction normal to the solid-liquid interface is favored. This is illustrated in Fig. 12. Grains with the preferred orientation will grow at the expense of adjacent grains that are less favorably oriented (Ref 8, 9, 10, 11, 12). Because dendrites follow the same preferred growth directions, those with the proper orientation will be favored. The requirement for a constant weld pool size becomes somewhat more complex when the effect of crystallographic orientation is considered (Fig. 12). Equation 14 can be modified to become:

R' = V COS φ /COS ( φ ' - φ )

(EQ 15)

where R' is the growth rate in the preferred direction required to maintain a constant weld pool shape and φ ' is the angle between R' and V, and ( φ ' - φ ) is the angle between the local preferred growth direction and R, which is normal to the solid-liquid interface.

FIG. 12 SCHEMATIC SHOWING MOVEMENT OF A CURVED SOLID-LIQUID INTERFACE FOR SEVERAL GRAINS (A AND B), AND THE CHANGE IN THE RAPID GROWTH DIRECTION RELATIVE TO THE INTERFACE POSITION.

Figure 12 shows the situation at successive positions of the weld pool as it moves in the direction of V. The situation for the weld moving from position 1 to position 2 is shown. The weld pool is assumed to be elliptical in shape, so that at any point (except at the very back of the weld pool where φ = 0) as the pool moves, the angle between R (or R') and V changes. If an adjacent grain becomes oriented such that R' becomes normal to the solid-liquid interface [that is, so that ( φ ' - φ ) is equal to zero], it will become most favored and will grow at the expense of adjacent grains. In Fig. 12, between positions 0 and 1, grain A is more favorably oriented than grain B. The reverse is true between positions 1 and 2. Figure 12 uses only a few widely spaced positions, but in reality there is a continuum of such positions.

Competitive growth causes grains to try to rotate into the direction of V. Figure 13(a) shows how the grains rotate in an elliptical weld pool. Dendrites will also try to rotate into the R direction, although the grains or dendrites do not actually rotate. Rather, the growth shifts to the most favored (or ) direction of the six present in the fcc, bcc, or hcp lattices. As grains are edged out by more favorably oriented grains, they get smaller and eventually can disappear. The solute that they were rejecting is spread out as this competitive process favors and disfavors different grains. Those grains or dendrites that eventually grow into the V direction will reject their solute into the weld pool, which gradually becomes solute enriched, contributing to possible crater cracking when the weld terminates. To combat this, the weld power should be gradually decreased in order to maintain a steep temperature gradient (minimizing the degree of constitutional supercooling). The use of a runoff weld tab could also be considered.

FIG. 13 SCHEMATIC SHOWING EFFECT OF HEAT INPUT AND WELDING SPEED VARIATIONS ON WELD GRAIN STRUCTURE. (A) LOW HEAT INPUT AND LOW WELDING SPEED, PRODUCING AN ELLIPTICAL WELD POOL. (B) HIGH HEAT INPUT AND HIGH WELDING SPEED, PRODUCING A TEAR-DROP-SHAPED WELD POOL. HERE, THE HEAT INPUT AND WELDING SPEED ARE NOT YET SUFFICIENT TO CAUSE HETEROGENEOUS NUCLEATION AT THE WELD POOL CENTERLINE. (C) HIGH HEAT INPUT AND HIGH WELDING SPEED, WITH HETEROGENEOUS NUCLEATION AT THE WELD POOL CENTERLINE. SOURCE: REF 13

The second factor to consider is the ability to extract heat at the solid-liquid interface. Heat is extracted by the temperature gradient of the solid. This gradient is a minimum at the back of the weld pool, where the temperature of the solid is the highest. Unfortunately, this is exactly where the solidification rate must be a maximum to keep up with V. The temperature gradient is a maximum at the side of the weld pool, where the least heat extraction is required, because here R = 0. As long as the maximum possible solidification rate, Rmax, is equal to or greate than V, the weld pool can maintain an elliptical shape. If V is increased above Rmax, then the solidification rate at the back of the weld pool cannot keep up, and

the weld pool elongates. Increasing V decreases the heat input, which shrinks the weld pool volume. With the weld elongating, this means that the diameter must decrease as the pool elongates. This causes the weld to take on a tear-drop shape, as shown in Fig. 13(b). The angle at the back of the weld pool will decrease until φ critical is reached. φ critical is given by: φ CRITICAL

= COS-1 (RMAX/V)

(EQ 16)

Rmax is controlled by the thermal conductivity of the solid, the geometry of the part being welded (primarily the thickness), and the preferred direction growth rate. When there is a shift to a tear-drop weld pool, the grain (or dendrite) rotation is limited to φ critical. When this situation is attained (Fig. 13b), the grains (or dendrites) will all grow together at the center of the weld pool. The solute that is being rejected into the terminal transient will build up at the centerline, which can lead to centerline cracking. This cracking is caused by the lower solidification temperature of this solute-rick centerline region. As the surrounding material solidifies, stresses are developed that cannot be supported by the liquid centerline material, causing centerline cracking. The obvious remedy is to reduce the welding rate, although this has a production rate penalty. Figure 13(b) typifies what is to be expected from a pure material or one with a low solute content. In the case of alloys or in cases when the solute content is high, Fig. 8 predicts that when R is increased, not only does the weld pool become tear-drop shaped, but the nature of the structure changes, as shown in Fig. 11. If R is large enough, equiaxed grains can develop from heterogenous nucleation in the supercooled centerline material (Fig. 13c). Such behavior has indeed been observed in aluminum alloys (Ref 13, 14, 15). Figure 14 shows actual weld microstructures corresponding to the schematic structures of Fig. 13.

FIG. 14 ACTUAL WELD MICROSTRUCTURES CORRESPONDING TO SCHEMATICS OF FIG. 13. (A) GRAIN STRUCTURE OF A GTAW OF ALUMINUM ALLOY 6061, MADE WITH Q OF 700 W (200 BTU/H) AND V OF 5.1 MM/S (0.20 IN./S). (B) GRAIN STRUCTURE OF A GTAW OF PURE ALUMINUM, MADE WITH V OF 20.8 MM/S (0.819 IN./S). (C) GRAIN STRUCTURE OF A GTAW OF ALUMINUM ALLOY 6061, MADE WITH A HIGHER Q THAN THAT OF (A) (1320 W, OR 385 BTU/H) AND V OF 12.7 MM/S (0.500 IN./S). GTAW, GAS-TUNGSTEN ARC WELD. SOURCE: REF 13, 15


7. F. MATSUDA, T. HASHIMOTO, AND T. SENDA, TRANS. NATL. RES. INST. MET. (JPN), VOL 11, 1969, P 43 8. W.F. SAVAGE, C.D. LUNDIN, AND A.H. ARONSON, WELD. J., VOL 44, 1965, P 420S-425S

9. W.F. SAVAGE, WELD. WORLD, VOL 18, 1980, P 89-113 10. W.F. SAVAGE, C.D. LUNDIN, AND R.J. HRUBEC, WELD. J., VOL 47, 1968, P 420S-425S 11. W.F. SAVAGE, C.D. LUNDIN, AND T.F. CHASE, WELD. J., VOL 47, 1968, P 522S-526S 12. W.F. SAVAGE, E.F. NIPPES, AND J.S. ERICKSON, WELD. J., VOL 55, 1976, P 213S-221S 13. S. KOU AND L. LE, METALL. TRANS. A, VOL 19, 1988, P 1075-1082 14. T. GANAHA, B.P. PEARCE, AND H.W. KERR, METALL. TRANS. A, VOL 11, 1980, P 1351-1359 15. H. NAKAGAWA, M. KATOH, F. MATSUDA, AND T. SENDA, TRANS. JPN. WELD. SOC., VOL 4, 1973, P 11 Fundamentals of Weld Solidification Harvey D. Solomon, General Electric Company

Nonequilibrium Effects: High-Rate Weld Solidification and Composition Banding As has been noted, while the constitutional supercooling model (described in previous sections of this article) gives valuable insights into the establishment of nonplanar structures, it does not necessarily give an exact depiction of the composition gradient develop within the dendrites or cells that may form and within the interdendritic or intercellular regions. This has been shown by Brooks and Baskes (Ref 16) in their study of gas-tungsten arc and electron-beam aluminum-copper and iron-niobium welds. For the aluminum-copper welds, the cell core compositions were about 2 to 3 times the KC0 value expected for the first material to solidify (see Fig. 5) and almost 4 times KC0 for the iron-niobium welds. This discrepancy was ascribed to the undercooling because of the curvature at the cell tip (see Eq 4), which is not considered in the constitutional supercooling model. Tip undercooling decreases the temperature at which the solidification occurs and, as Fig. 3 shows, increases the composition of the solid that is forming. No initial transient of the sort shown in Fig. 5 was observed. This was explained by using the lower solidification rate at the side of a cell, rather than that at the dendrite tip in the calculation of the size of the initial transient. This increases the length of the transient (which is inversely proportional to the solidification rate, R) and eliminates the sharp composition variation of the initial transient shown in Fig. 5. Brooks and Baskes also incorporated postsolidification solid-state diffusion of solute (following the approach of Brody and Flemings, which is given in Ref 17) to explain the composition gradients that they measured. Such diffusion is also not considered in the conventional constitutional supercooling model. In addition, the classical theory of constitutional supercooling, as described by Eq 10, 11, and 12, does not take into account the temperature gradient that develops in the solid, the thermal conductivity of the solid and the liquid, and, most importantly, the surface tension of the solid-liquid interface. These factors are considered in interface stability models (Ref 18, 19, 20, 21, 22, 23, 24, 25, 26), which describe the breakdown of a planar solid-liquid interface in terms of the stability of nonplanar perturbations. The importance of these models is that they explain the observed transition from nonplanar solidification back to planar solidification when the solidification rate is increased to very high levels. As has already been shown, at low solidification rates there can be a transition from planar to nonplanar solidification, as described by constitutional supercooling (as shown in Fig. 8). As the solidification rate is increased, the radius of curvature and spacing of the dendrite arms will decrease, as does the cell spacing. What eventually occurs, at very high solidification rates, is that the very small radius that is required is not supportable, because of the increased surface energy. The stability models predict that there will be a solidification rate (the limiting stability rate) beyond which there will be a transition back to planar solidification. Figure 15 shows the calculated dendrite tip radius as a function of velocity of the solid-liquid interface. The increase in R at high rates defines the limiting stability rate.

FIG. 15 PLOT OF DENDRITE TIP RADIUS, R, VERSUS VELOCITY, V, OF THE SOLID-LIQUID INTERFACE FOR AN AG-5% CU COMPOSITION WITH TEMPERATURE GRADIENT, G, OF 105 K/CM (5 × 105 °F/IN). CURVE A HAS DL(T) AND K(V) WITH KE AND ML AS CONSTANTS; CURVE B, DL (T) WITH KE AND ML AS CONSTANTS; CURVE C, WITH DL, KE, AND ML AS CONSTANTS. THE VELOCITY WHERE THE TIP RADIUS INCREASES IS THE STABILITY LIMIT. THE CALCULATIONS ARE MADE USING THE CONVENTIONAL STABILITY LIMIT APPROACH, WITH THE ADDITION OF A TEMPERATURE-DEPENDENT DIFFUSION COEFFICIENT, AND WITH THE ADDITION OF BOTH A TEMPERATURE-DEPENDENT DIFFUSION COEFFICIENT AND A VELOCITY-DEPENDENT PARTITION COEFFICIENT. SOURCE: REF 21

Each of the three curves in Fig. 15 illustrates different assumptions. The right-most curve shows the result using the approach of Mullins and Sekerka (Ref 18), who developed the original interface stability model. Here, the solute diffusion rate in the liquid, DL, the equilibrium distribution coefficient, K, and the liquidus slope, ML, are all assumed to be constants. The middle curve takes into account the fact that the diffusion coefficient is temperature dependent. (Increasing the degree of undercooling that is due to rapid solidification decreases both the solute diffusion rate and the limiting velocity.) The left-most curve not only considers the temperature dependence of the diffusion coefficient, but also takes into account a dependency of the distribution coefficient on the solidification rate. The preceding discussions have utilized the distribution coefficient, K, determined from equilibrium phase diagrams. At high solidification rates, this equilibrium is not achieved at the solid-liquid interface. Rather, K is a function of the solidification rate. One of several different expressions for the variation of K with V is given by (Ref 23):

K = [KE + (V A0/DL)] / [1 + (VA0/DL)]

(EQ 17)

where Ke is the equilibrium distribution coefficient (K in Eq 8 and 12), and a0 is a distance that is related to the interatomic spacing. Different models use different definitions for a0 and different functional relationships, but give similar results. When V > D/a0, K will approach 1. For D = 5 × 10-5 cm2/s, or 5 × 10-8 ft2/s (a typical value for liquid diffusion), and a0 = 5 × 10-8 cm, or 2.0 × 10-8 in., K will approach 1 for V > 103 cm/s, or 4 × 102 in./s. With K = 1, there is no partition of the solute, no constitutional supercooling, and the solidification will be planar for a positive temperature gradient. Figure 15 shows that the limiting velocity, as predicted by the stability model, is quite a bit lower than 103 cm/s (4 × 102 in./s). Thus, the reestablishment of planar solidification is not just the result of K approaching unity. Rather, it is due to the lack of the stability of nonplanar protrusions, which is influenced by this change in K. This development of planar solidification is important because it can arise in electron-beam and laser welds, and because the development of planar solidification lessens the tendency for hot cracking. Figure 15 shows the results of calculations, whereas Fig. 16 shows the results of actual observations made on silvercopper electron-beam scans (passes of the electron beam). These results are in general agreement with the predictions of Fig. 15. Figure 16 shows the compositional dependence of the transition to segregation-free planar solidification. It is

analogous to Fig. 8, which describes the transition from planar to nonplanar solidification that occurs at lower solidification rates. Figure 16 also shows that just below the limiting solidification rate, a banded structure can develop. This banding develops at a constant solidification rate and is due to instabilities at the solid-liquid interface.

FIG. 16 OBSERVED MICROSTRUCTURES IN SILVER-COPPER ALLOYS AS A FUNCTION OF THE SCAN VELOCITY (BEAM TRANSLATION VELOCITY) FOR ELECTRON-BEAM WELDS. OPEN CIRCLES DENOTE MICROSEGREGATIONFREE STRUCTURE; CLOSED CIRCLES, A STRUCTURE CONSISTING OF CELLS OR DENDRITES; AND OPEN TRIANGLES, THE DEVELOPMENT OF A BANDED STRUCTURE. SOURCE: REF 24

A more common type of banding is associated with the ripples observed on the surface of most welds. Variations in the welding power will cause variations in the solidification rate. If the power decreases, then the solidification rate will increase. More solute will be incorporated into the solid and a solute-rich band will be created (for K < 1 and where ML is negative). The reverse is true for an increase in the power (that is, the solidification rate will decrease), and this will cause a solute-poor band to form (for K < 1 and where ML is negative). These power fluctuations can be deliberate, as in pulsed welding, or can result from unavoidable instabilities in the welding power supplies or fluctuations in primary voltage, mechanical (travel) motion, and so on. Solute banding and surface rippling are illustrated in Fig. 17.

FIG. 17 SOLUTE BANDING AND SURFACE RIPPLING IN WELD BEADS. (A) SURFACE RIPPLES IN A TUNGSTENINERT GAS WELD MADE WITH AL-5SI FILLER. 6×. (B) SULFUR AND PHOSPHORUS SOLUTE BANDING AT THE FUSION BOUNDARY OF A MANUAL ARC WELD IN CARBON-MANGANESE STEEL. 26×. COURTESY OF THE WELDING INSTITUTE

Variations in the welding speed can also cause banding, as can the waves associated with the weld pool circulation, which can also influence the solidification rate. Special alternating magnetic fields have been used to supplement the surface tension forces, the usual electromagnetic forces, and the arc pressure and buoyancy forces, all of which affect the weld pool circulation (Ref 27). The aim is to induce interface fragmentation and cause grain refinement. The solidified structure can also be influenced by several other factors, in addition to variations in the rate of movement of the solid-liquid interface (Ref 27). Inoculants have been used to nucleate equiaxed grains by heterogeneous nucleation on the inoculant particles. Ferro-titanium, ferro-niobium, titanium carbide, and titanium nitride have been used as grain refiners in steels. Zirconium and titanium have also been used as inoculants for aluminum. The effectiveness of inoculants depends not only on the type of inoculants used, but on the size of the inoculant particles, the method of introduction (where in the weld pool it is added), and welding conditions. Arc motion and vibrations have also been used to suppress columnar growth and to produce a finer grain size (Ref 27). The effectiveness of arc vibration on the as-solidified grain size of an Al-2.5% Mg weld is illustrated in Fig. 18. The effectiveness of this approach depends on the frequency and amplitude of the arc motion, the heat input, arc length, and sheet thickness. In general, the greater the amplitude of vibration, the less the frequency required for grain refinement (Ref 28).

FIG. 18 PLOT OF GRAIN SIZE VERSUS ARC VIBRATION AMPLITUDE FOR AL-2.5MG WELD. WELD PARAMETERS: WELDING CURRENT, 150 A; WELDING SPEED, 200 MM/MIN (50 IN./MIN); ARC GAP, 2.4 MM (0.094 IN.); ARC VIBRATION FREQUENCY, 10 HZ. SOURCE: REF 27

Campbell (Ref 28) has developed frequency-amplitude maps that define several physical effects associated with the use of vibrations (that is, peak pressures generated, development of shearing stresses, power densities, laminar flow around dendrite arms, and the development of standing waves). These factors affect microstructural features, such as grain size, amount of microsegregation, and the development of porosity. For a given frequency of vibration, increasing the amplitude of vibration can cause a refinement of the grain size (Fig. 18), reduce microsegregation, and, unfortunately, increase the degree of porosity that is caused by vibration-induced cavitation. Vibrations can also cause splatter of the liquid weld metal; coarsen intragranular features, such as the secondary dendrite arm spacing; and promote hot tearing. Great care must therefore be taken when using vibratory motion and when determining the optimum amplitude and frequency of vibration. Variations in the arc current can also help to refine the weld microstructure. Ultrasonic vibration of the weld pool has also been found to be effective in some cases. These vibrations have been added through the filler wire and through the use of water-cooled copper probes (in large weld pools, such as those developed in electro-slag welding). In some cases, variations in the nature of the weld pool circulation have also been found to be effective in refining the weld structure. The desire to produce a fine-grained equiaxed weld structure is driven by the improved mechanical properties and lessened hot cracking that this structure provides. Torch weaving (Ref 29) (that is, moving the arc in a sinusoidal rather than linear fashion) has been used to break up the linear nature of the weld and to prevent hot cracks from propagating along the weld. A weld made with torch weaving is shown in Fig. 19. This weaved weld structure prevents cracks from running the length of the weld, but adds to the cost of the weld by decreasing the linear travel speed.

FIG. 19 TYPICAL GRAIN STRUCTURE OBTAINED WHEN TORCH WEAVING IS USED TO WELD ALUMINUM ALLOY 2014. WELD PARAMETERS: TRANSVERSE ARC FREQUENCY OSCILLATION, 1.0 HZ; AMPLITUDE OF OSCILLATION, 1.5 MM (0.059 IN.). SOURCE: REF 13, 30

Our understanding of weld solidification owes much to the more extensive study of the solidification of castings. The reader should refer to that subject area for a more in-depth treatment of solidification in general. However, it should be remembered that the boundary conditions for weld solidification and for that of a casting are not the same. Therefore, care must be taken in applying all that has been learned from casting metallurgy to the metallurgy of welds.


13. S. KOU AND L. LE, METALL. TRANS. A, VOL 19, 1988, P 1075-1082 16. J.A. BROOKS AND M.I. BASKES, WELD MICROSTRUCTURE CHARACTERIZATION AND MODELING, RECENT TRENDS IN WELDING SCIENCE AND TECHNOLOGY, S.A. DAVID AND J.M. VITEK, ED., ASM INTERNATIONAL, 1990, P 93-99 17. H.D. BRODY AND M.C. FLEMINGS, TRANS. AIME, VOL 236, 1966, P 615-624 18. W.W. MULLINS AND R.F. SEKERKA, J. APPL. PHYS., VOL 35, 1964, P 444-451 19. J.S. LANGER AND H. MULLER-KRUMBHAAR, ACTA METALL., VOL 26, 1978, P 1681-1687 20. W. KURZ AND D.J. FISCHER, ACTA METALL., VOL 29, 1981, P 11-20 21. W. KURZ, B. GIOVANOLA, AND R. TRIVEDI, ACTA METALL., VOL 34, 1986, P 823-830 22. R. TRIVEDI AND W. KURZ, ACTA METALL., VOL 34, 1986, P 1663-1670 23. M.J. AZIZ, J. APPL. PHYS., VOL 53, 1982, P 1158-1168 24. W.J. BOETTINGER, D. SHECHTMAN, R.J. SCHAEFFER, AND BIANCANIELLO, METALL. TRANS. A, VOL 15, 1984, P 55-66 25. W.J. BOETTINGER AND S.R. CORIELL, MATER. SCI. ENG., VOL 65, 1984, P 27-36 26. M. CARRARD, M. GREMAUD, M. ZIMMERMANN, AND W. KURZ, ACTA METALL., VOL 40, 1992, P 983-996 27. G.J. DAVIES AND J.G. GARLAND, INT. MET. REV., VOL 20, 1975, P 83-106 28. J. CAMPBELL, INT. MET. REV., VOL 26 (NO. 2), 1981, P 71-108 29. S. KOU, RECENT TRENDS IN WELDING SCIENCE AND TECHNOLOGY, S.A. DAVID AND J.M. VITEK, ED., ASM INTERNATIONAL, 1989, P 137-147 30. S. KOU, WELDING METALLURGY, JOHN WILEY & SONS, 1987, P 161-162 Fundamentals of Weld Solidification Harvey D. Solomon, General Electric Company

References

1. J.D VERHOEVEN, FUNDAMENTALS OF PHYSICAL METALLURGY, JOHN WILEY & SONS, 1975 2. H.D. SOLOMON, IN TREATISE ON MATERIALS SCIENCE AND TECHNOLOGY, VOL 25, EMBRITTLEMENT OF ENGINEERING ALLOYS, C.L. BRIANT AND S.K. BANERJI, ED., ACADEMIC PRESS INC., 1983, P 525-599 3. J.A. BURTON, R.C. PRIM, AND J. SLICHTER, J. CHEM. PHYS., VOL 21, 1953, P 1987-1991 4. W.A. TILLER, K.A. JACKSON, J.W. RUTTER, AND B. CHALMERS, ACTA METALL., VOL 1, 1953, P 428-437 5. W.A. TILLER AND J.W. RUTTER, CAN. J. PHYS., VOL 34, 1956, P 96-121 6. W.F. SAVAGE, E.F. NIPPES, AND T.W. MILLER, WELD. J., VOL 55, 1976, P 165S-173S 7. F. MATSUDA, T. HASHIMOTO, AND T. SENDA, TRANS. NATL. RES. INST. MET. (JPN), VOL 11, 1969, P 43 8. W.F. SAVAGE, C.D. LUNDIN, AND A.H. ARONSON, WELD. J., VOL 44, 1965, P 420S-425S 9. W.F. SAVAGE, WELD. WORLD, VOL 18, 1980, P 89-113 10. W.F. SAVAGE, C.D. LUNDIN, AND R.J. HRUBEC, WELD. J., VOL 47, 1968, P 420S-425S 11. W.F. SAVAGE, C.D. LUNDIN, AND T.F. CHASE, WELD. J., VOL 47, 1968, P 522S-526S 12. W.F. SAVAGE, E.F. NIPPES, AND J.S. ERICKSON, WELD. J., VOL 55, 1976, P 213S-221S 13. S. KOU AND L. LE, METALL. TRANS. A, VOL 19, 1988, P 1075-1082 14. T. GANAHA, B.P. PEARCE, AND H.W. KERR, METALL. TRANS. A, VOL 11, 1980, P 1351-1359 15. H. NAKAGAWA, M. KATOH, F. MATSUDA, AND T. SENDA, TRANS. JPN. WELD. SOC., VOL 4, 1973, P 11 16. J.A. BROOKS AND M.I. BASKES, WELD MICROSTRUCTURE CHARACTERIZATION AND MODELING, RECENT TRENDS IN WELDING SCIENCE AND TECHNOLOGY, S.A. DAVID AND J.M. VITEK, ED., ASM INTERNATIONAL, 1990, P 93-99 17. H.D. BRODY AND M.C. FLEMINGS, TRANS. AIME, VOL 236, 1966, P 615-624 18. W.W. MULLINS AND R.F. SEKERKA, J. APPL. PHYS., VOL 35, 1964, P 444-451 19. J.S. LANGER AND H. MULLER-KRUMBHAAR, ACTA METALL., VOL 26, 1978, P 1681-1687 20. W. KURZ AND D.J. FISCHER, ACTA METALL., VOL 29, 1981, P 11-20 21. W. KURZ, B. GIOVANOLA, AND R. TRIVEDI, ACTA METALL., VOL 34, 1986, P 823-830 22. R. TRIVEDI AND W. KURZ, ACTA METALL., VOL 34, 1986, P 1663-1670 23. M.J. AZIZ, J. APPL. PHYS., VOL 53, 1982, P 1158-1168 24. W.J. BOETTINGER, D. SHECHTMAN, R.J. SCHAEFFER, AND BIANCANIELLO, METALL. TRANS. A, VOL 15, 1984, P 55-66 25. W.J. BOETTINGER AND S.R. CORIELL, MATER. SCI. ENG., VOL 65, 1984, P 27-36 26. M. CARRARD, M. GREMAUD, M. ZIMMERMANN, AND W. KURZ, ACTA METALL., VOL 40, 1992, P 983-996 27. G.J. DAVIES AND J.G. GARLAND, INT. MET. REV., VOL 20, 1975, P 83-106 28. J. CAMPBELL, INT. MET. REV., VOL 26 (NO. 2), 1981, P 71-108 29. S. KOU, RECENT TRENDS IN WELDING SCIENCE AND TECHNOLOGY, S.A. DAVID AND J.M. VITEK, ED., ASM INTERNATIONAL, 1989, P 137-147 30. S. KOU, WELDING METALLURGY, JOHN WILEY & SONS, 1987, P 161-162 Fundamentals of Weld Solidification Harvey D. Solomon, General Electric Company

Selected References

• CASTING, VOL 15, ASM HANDBOOK, ASM INTERNATIONAL, 1988, P 100-181

• S.A. DAVID AND J.M. VITEK, INT. MATER. REV., VOL 34, 1989, P 213-245 • J.H. DEVLETION AND W.E. WOOD, VOL 6, METALS HANDBOOK, 9TH ED., AMERICAN SOCIETY FOR METALS, 1983, P 21-49 • K. EASTERLING, INTRODUCTION TO THE PHYSICAL METALLURGY OF WELDING, BUTTERWORTHS, 1985 • M.C. FLEMINGS, SOLIDIFICATION PROCESSING, MCGRAW-HILL, 1974 • W. KURZ AND D.J. FISHER, FUNDAMENTALS OF SOLIDIFICATION, 3RD ED., TRANS TECH, 1989 • J.F. LANCASTER, METALLURGY OF WELDING, 3RD ED., GEORGE ALLEN & UNWIN, 1980 • M. RAPPAZ, INT. MATER. REV., VOL 34, 1989, P 93-123 Nature and Behavior of Fluxes Used for Welding D.L. Olson, S. Liu, R.H. Frost, G.R. Edwards, and D.A. Fleming, Colorado School of Mines

Introduction FLUXES are added to the welding environment to improve arc stability, to provide a slag, to add alloying elements, and to refine the weld pool (Ref 1, 2). Different ingredients in the flux system will provide the process with different pyrometallurgical characteristics and thus different weld metal properties (Ref 3, 4). The slag that forms during welding covers the hot weld metal and protects it from the atmosphere. Welding slag consists of the glass-forming components of the flux, as well as inclusions that form in the weld pool, coalesce, rise, and become incorporated into the slag. The need to improve flux formulation to achieve optimal weld metal composition, and ultimately improve the properties of weldments, has led to fundamental studies of weld pyrochemistry. Understanding the thermodynamic and kinetic factors that are prevalent at the electrode, in the arc column, and in the weld pool, has led to more precise prediction of the final weld metal composition (Ref 5, 6, 7, 8, 9, 10, 11, 12, 13, 14, 15).

References

1. G.E. LINNERT, CHAPTER 8, WELDING METALLURGY, VOL 1, AWS, 1965, P 367-396 2. C.E. JACKSON, FLUXES AND SLAGS IN WELDING, WELD. RES. BULL., NO. 190, 1973 3. T. LAU, G.C. WEATHERLY, AND A. MCLEAN, THE SOURCES OF OXYGEN AND NITROGEN CONTAMINATION IN SUBMERGED ARC WELDING USING CAO-AL2O3 BASED FLUXES, WELD. J., VOL 64 (NO. 12), 1985, P 343S-347S 4. T.H. NORTH, H.B. BELL, A. NOWICKI, AND I. CRAIG, SLAG/METAL INTERACTION, OXYGEN, AND TOUGHNESS IN SUBMERGED ARC WELDING, WELD. J., VOL 57 (NO. 3), 1978, P 63S-75S 5. N. CHRISTENSEN AND J. CHIPMAN, SLAG-METAL INTERACTION IN ARC WELDING, WELD. RES. BULL., NO. 15, JAN 1953, P 1-14 6. R.H. FROST, D.L. OLSON, AND S. LIU, PYROCHEMICAL EVALUATION OF WELD METAL INCLUSION EVOLUTION, PROC. 3RD INT. CONF. TRENDS IN WELDING, ASM INTERNATIONAL, JUNE 1992 7. C.A. NATALIE, D.L. OLSON, AND M. BLANDER, PHYSICAL AND CHEMICAL BEHAVIOR OF WELDING FLUXES, ANN. REV. MATER. SCI., VOL 16, 1986, P 389-413 8. J.E. INDACOCHEA, M. BLANDER, N. CHRISTENSEN, AND D.L. OLSON, CHEMICAL REACTIONS WITH FEO-MNO-SIO2 FLUXES, METALL. TRANS. B, VOL 16, 1985, P 237-245 9. U. MITRA AND T.W. EAGAR, SLAG-METAL REACTIONS DURING WELDING, METALL. TRANS. B, VOL 22, 1991, P 65-100

10. C.S. CHAI AND T.W. EAGAR, SLAG-METAL EQUILIBRIUM DURING SUBMERGED ARC WELDING,METALL. TRANS. B, VOL 12, 1981, P 539-547 11. N. CHRISTENSEN, METALLURGICAL ASPECTS OF ARC WELDING, WELD. J., VOL 27, 1949, P 373S-380S 12. O. GRONG, D.L. OLSON, AND N. CHRISTENSEN, CARBON OXIDATION IN HYPERBARIC MMA WELDING, MET. CONSTRUCT., VOL 17. DEC 1985, P 810R-814R 13. U. MITRA AND T.W. EAGAR, SLAG-METAL REACTIONS DURING SUBMERGED ARC WELDING OF ALLOY STEELS, METALL. TRANS. A, VOL 15, 1984, P 217-227 14. T.W. EAGAR, SOURCES OF WELD METAL OXYGEN CONTAMINATION DURING SUBMERGED ARC WELDING, WELD. J.,VOL 57, 1978, P 76S-80S 15. N. CHRISTENSEN, WELDING METALLURGY, LECTURE NOTES, NTH, 1979 Nature and Behavior of Fluxes Used for Welding D.L. Olson, S. Liu, R.H. Frost, G.R. Edwards, and D.A. Fleming, Colorado School of Mines

Equilibrium Parameters Equilibrium is not achieved during welding, because of the very large temperature and density gradients, the short reaction times, and the large electric currents. Despite these expected departures from equilibrium, thermodynamic considerations can be used as a guide for constraining chemical reactions and mechanisms involved in welding (Ref 5, 7, 9). A common approach is to assume that extremely high temperatures and high surface-to-volume ratios allow thermodynamic equilibrium to be locally attained, in spite of the short reaction times available. Further complicating the issue is the dependence on welding parameters of the chemical partitioning between the slag and the weld metal (Ref 3, 7). This dependence suggests that the pyrometallurgical reactions involved are influenced by the processes that occur at the electrode tip. As a specific example, the welding parameters affect metal-droplet size, which in turn alters the chemical kinetics. Nonetheless, thermodynamics will reveal the direction taken by the chemical reactions (but will not accurately predict weld metal composition). The effective slag-metal reaction temperature has been estimated (by pyrochemical analysis of slag-metal compositions) to be approximately 1900 °C (3450 °F), a temperature intermediate between that of the hot spot at the arc root (2300 °C or 4170 °F) and the melting point of iron (1500 °C or 2730 °F). Thus, during the droplet lifetime, the average temperature is effectively 1900 °C (2730 °F). During this period, oxygen from the hot spot reactions is distributed throughout the droplet, and that oxygen reacts with the metallic elements to form oxides; these products of oxidation pass into the slag, and the slag-metal reactions tend toward equilibrium. Estimates of the time for which the molten slag and molten metal are in contact range from 3 to 8 s. During the process, the gaseous phase in the arc cavity contacts the metal for an estimated 0.5 to 1.0 s. Effect of Oxygen The most important chemical reagent in controlling weld metal composition, and thus microstructure and properties, is oxygen (Ref 3, 4, 14, 16, 17, 18, 19). Oxygen directly reacts with alloying elements to alter their effective role by: • • •

REDUCING HARDENABILITY PROMOTING POROSITY PRODUCING INCLUSIONS.

All three effects are significant to weld quality. Oxygen is introduced into the weld pool at high temperatures by:

• • • •

THE PRESENCE OF OXIDE FLUXES THAT DISSOCIATE IN THE ARC THE SLAG-METAL REACTIONS IN THE WELD POOL THE OXIDES ON THE SURFACE OF BAKED METALLIC POWDERS MIXED WITH FLUX OR ON ELECTRODE THE ASPIRATION OF ATMOSPHERE (AIR) INTO THE ARC.

Shielding gas may contain oxidizing reagents; a common gas is 75Ar-25CO2 or 100% CO2. As a 100 to 1000 ppm, depending on the type of welding consumable used. The weld metal oxygen measured directly at the molten electrode tip has been reported to be as high as 1400 ppm. Individual droplets have been found to contain as much as 2000 ppm O2. There are two views concerning the genesis of the high oxygen content. One suggests that pyrochemical or electrochemical reactions (or both) provide oxygen to the electrode tip, and then further oxidation occurs within the droplet as it passes through the arc. The other view is that the high oxygen levels of the droplet represent the maximum buildup of oxygen in the electrode prior to detachment, with limited reaction during flight across the arc to the weld pool. In either case, high oxygen concentrations are introduced into the weld pool by the welding process. At the melting point, the solubility of oxygen in pure liquid iron is approximately 1600 ppm at 100 kPa (1 atm) pressure. During solidification, solubility decreases to about 860 ppm at 1500 °C (2730 °F) in δ-Fe. Most of the alloying elements present in liquid steel reduce oxygen solubility through deoxidation equilibria. Steelmaking processes typically yield analytical oxygen levels ranging from 70 to 100 ppm. Welds typically pick up oxygen to levels of several hundred ppm, then deoxidize to oxygen levels of around 200 to 300 ppm. Deoxidation of the weld metal occurs in two separate steps, the first being the primary deoxidation of the weld pool (Ref 6). Secondary deoxidation occurs during solidification as solute concentrations increase within the intercellular or interdendritic regions. The secondary deoxidation will either form very small inclusions or will coat the interdendritically trapped primary inclusions. The high oxygen concentrations added to the weld pool by the metal droplets significantly affect deoxidation. Figure 1 shows experimentally measured soluble oxygen concentrations for various deoxidants (solid lines), along with deoxidation curves predicted by the solubility products (broken lines). The experimental deviation is caused by interactions of the deoxidant with other alloy elements. If the oxygen and metallic element concentrations resulting from the welding process exceed the equilibrium concentration for a specific reaction, inclusions will result. The ability to form a specific inclusion will correspond directly to the position of the weld pool composition relative to the activity plot for this inclusion. Thus, the thermodynamic order for the formation of primary oxides would be: Al2O3 > Ti2O3 > SiO2 > Cr2O4 > MnO.

FIG. 1 DEOXIDATION EQUILIBRIA IN LIQUID IRON ALLOYS AT 1600 °C (2910 °F). THE BROKEN LINES SHOW DEOXIDATION EQUILIBRIA PREDICTED BY SOLUBILITY PRODUCT CALCULATIONS. THE SOLID LINES SHOW EXPERIMENTALLY DETERMINED SOLUBLE OXYGEN CONCENTRATIONS FOR VARIOUS DEOXIDANTS. THE EXPERIMENTAL DEVIATION IS CAUSED BY VARIATIONS IN THE ACTIVITY COEFFICIENTS WITH INCREASING DEOXIDANT CONCENTRATION. SOURCE: REF 20

Inclusion Formation. Inclusions form as a result of reactions between metallic alloy elements and nonmetallic tramp elements, or by mechanical entrapment of nonmetallic slag or refractory particles. Inclusions may include: • • • • • •

OXIDES SULFIDES NITRIDES CARBIDES OTHER COMPOUNDS MULTIPLE PHASES

Among these, oxides and complex oxides occur most frequently in the size range known to influence weld metal microstructure.

Using only thermodynamic considerations in the analysis of slag-metal reactions, the following reactions may be considered as the ones that describe inclusion formation:

X M + Y O = (MXOY)

(EQ 1)

where the underlining of a component M ( M ) and the component O ( O ) indicate that M and O, respectively, are dissolved in the metal, and:

X M + Y(FEO) = (MXOY) + YFE

(EQ 2)

The equilibrium constants (K1 and K2) for the above two reactions are: K1 =

( AM x Oy ) [aM ]x [ao ] y

(EQ 3)

where ai is the activity, a function of the concentration, for component i, and: K2 =

( AM x Oy ) [aM ]x [aFeO ] y

(EQ 4)

Equilibrium compositions for the weld deposit can be estimated with Eq 1, 2, 3, and 4 and can be used to predict trends for the weld pool pyrochemical reactions. The actual compositions may, however, differ from the calculated values due to alloying-element partitioning during cellular or dendritic solidification, commonly observed in steel weldments. During solidification of a weld metal, solute elements segregate to the liquid at the solid/liquid interface, and the solute concentrations can reach high levels in the interdendritic regions, as suggested in Fig. 2. Neglecting solid diffusion, the solute composition in the liquid at the solid/liquid interface can be modeled (Ref 6) by the nonequilibrium lever rule, or Scheil Equation:

CL = C0FLK-1

(EQ 5)

where C0 is the bulk concentration in the weld pool, CL is the solute concentration in the liquid at the interface, and k is the equilibrium partition ratio.

FIG. 2 SCHEMATIC SHOWING THE SOLID AND LIQUID COMPOSITION PROFILES MODELED BY EQ 5 BASED ON THE ASSUMPTION THAT THERE IS COMPLETE LIQUID DIFFUSION AND NO SOLID DIFFUSION. C0, THE INITIAL ALLOY COMPOSITION; K, THE PARTITION RATIO OF THE SOLID TO LIQUID COMPOSITIONS ON THE EQUILIBRIUM PHASE DIAGRAM; KC0, THE INITIAL COMPOSITION OF THE SOLID; AND CS AND CL, THE SOLID AND LIQUID COMPOSITIONS, RESPECTIVELY, AT THE SOLIDIFICATION INTERFACE. SOURCE: REF 4

The equilibrium partition ratio, k, controls the direction and the extent of segregation. For most alloy elements in steel, the partition ratio is less than one, and the element segregates to the interdendritic liquid. Consider the deoxidation equilibrium as represented by the dissolution reaction for a complex oxide:

MXNYOZ = X M + Y N + Z O

(EQ 6)

The free energy change associated with the dissolution of the MxNyOz can be written as: ∆G = ∆G° + RT ln

[ M ]x [ N ] y [O]Z [ M x N y Oz ]

(EQ 7)

where [M], [N], and [O] are the solute activities in the liquid; and x, y, and z are the stoichiometric constants from Eq 6. The value [MxNyOz] is the activity for the specific inclusion and can be assumed to have the value of one. The ratio of the activities of the reactants to that of the oxide can be termed the activity quotient (Q): Q=

[ M ]x [ N ] y [O]Z [ M x N y Oz ]

(EQ 8)

With Eq 5 representing the extent of segregation, the solute activities in the interdendritic liquid can be written as:

[M]=[

M][MO]

[N]=[

N][N0]

[O] = [

O][O0]

(EQ 9A) (EQ 9B) (EQ 9C)

In Eq 9a, 9b, and 9c, [MO], [NO], and [OO] are the bulk concentrations of M, N, and O in the melt; and [γM], [γN], and [γO] are the activity coefficients for the solutes. Substituting Eq 9a, 9b, and 9c into Eq 8 gives the free energy as a function of the remaining liquid fraction:

(EQ 10)

Equation 10 expresses the free energy driving force for oxide dissolution. The first term on the right side, ∆G, represents equilibrium conditions. The second term represents the departure from equilibrium caused by changes in reactant concentrations or by segregation during solidification. At equilibrium, the free energy driving force, ∆G, is zero, and the equilibrium concentrations can be found by equating the two terms on the right-hand side of Eq 10. At equilibrium, the activity quotient becomes the equilibrium constant, Keq:

(EQ 11) Inclusion precipitation is possible when the concentrations of oxygen and deoxidants exceed the equilibrium values for a particular oxide. This condition can be expressed with the ratio of the activity quotient to the equilibrium coefficient:

(EQ 12) A precipitation index less than unity indicates that the concentrations of oxygen and deoxidants are below the equilibrium value, and the precipitation of the inclusions will not occur. A value greater than 1.0 indicates that concentrations are sufficiently high for precipitation according to the methodology described above. The compounds Al2O3, Ti2O3, and SiO2 are some of the oxides that will form in a low-carbon low-alloy steel weld. Multiple reactions can occur, and different oxides may appear in the same weld when more than one of these deoxidizers is present. Metal Transferability During Pyrochemical Reactions. The final weld metal concentration for a particular element is

made up of contributions from the filler wire, flux, and base metal; however, losses caused by the welding process vary for each element. Delta Quantity. The nominal composition of each weld can be calculated considering just the dilution effect of the filler

wire and base metal. The extent of loss or transfer of a specific element can be evaluated by a quantity, which expresses the difference between the analytical and the nominal composition. These quantities, designated delta quantity in this article, indicate the effect of the flux on element transfer during welding. A positive delta quantity indicates an elemental transfer from the flux to the weld metal. A negative delta quantity suggests an elemental loss from the weld pool to the slag. A null delta quantity for a specific element suggests an apparent equilibrium condition, in which the flux and slag content for that element are the same. Flux compositions with null delta-quantity behavior have been used to make equilibrium calculations and thus achieve a better understanding of the chemical reactions involved in welding (Ref 8). Investigators have also quantified elemental transfer by measuring similar neutral points (null delta quantity) for various flux systems and have developed a thermodynamic model capable of predicting neutral points for some slag systems (Ref 10). Arc Stabilizers. Arc welding fluxes are compositionally more complex than fluxes used in other metallurgical processes,

such as steelmaking. Many of the additions to the flux are not designed to assist weld metal refinement. Some additions are present to promote arc stability, generate plasma and protective gases, control viscosity, support out-of-position welding, and promote slag detachability. The welding arc requires an inert or chemically reducing plasma and shielding gas that can be easily ionized. Additions must be made to the flux to achieve the necessary current-carrying capacity and to maintain arc stability. Specific additions will be necessary for the various modes of current (direct current, dc, or alternating current, ac) and polarity. Alkali metal, zirconium, and titanium additions to the arc affect the ionization process and the ease with which the welding is reinitiated (reinitiation is required 50 to 60 times per second with ac welding). These additions come to the welding flux as feldspar (alkali aluminum silicates), rutile, lithium carbonate, titanium aluminate, and potassium oxalate, and they play an especially important role when welding either in the dc electrode negative (DCEN) mode or in the alternate current (ac) mode. Arc stabilizers are also important in high-speed welding when the cathode and anode spots are less stable. These additions are part of the AWS classification of electrodes, as seen by the fourth digit (E XXX X ) in the classification standard for steel electrodes for shielded metal arc welding (SMAW) given in Table 1. The electrodes designed to perform with ac mode or in the DCEN mode current have been specified as containing high titania or alkali metal (potassium or sodium) additions. Special additions such as Li2O have been used to achieve multipurpose results-for example, reducing the viscosity while increasing the arc stability.

TABLE 1 AWS CLASSIFICATION OF SELECTED SMAW ELECTRODES FOR WELDING MILD AND LOWALLOY STEELS

ELECTRODE DESIGNATION EXX10 EXXX1 EXXX2 EXXX3 EXXX4

CURRENT AND POLARITY(A) DCEP AC, DCEP AC, DCEN AC, DCEP, DCEN AC, DCEP, DCEN

PENETRATION ARC STABILIZERS DEEP DEEP MEDIUM LIGHT LIGHT

HIGH CELLULOSE-NA HIGH CELLULOSE-K HIGH RUTILE-NA HIGH RUTILE-K RUTILE-IRON POWDER

EXX24 EXXX5 EXXX6 EXX27

AC, DCEP, DCEN DCEP AC, DCEP AC, DCEN, DCEP

LIGHT MEDIUM MEDIUM MEDIUM

EXXX8

AC, DCEP

MEDIUM

EXX28

AC, DCEP

MEDIUM

RUTILE-IRON POWDER LOW H-NA LOW H-K IRON OXIDE-IRON POWDER LOW HYDROGEN-IRON POWDER LOW HYDROGEN-IRON POWDER

% FE 0-10 0 0-10 0-10 2540 50 0 0 50 2540 50

(A) DCEP, DC ELECTRODE POSITIVE, DCEN, DC ELECTRODE NEGATIVE Changes in Flux Composition With Delta Quantity. Figure 3 shows a transfer of manganese for welds made with

SiO2-TiO2-CaO-1Na2O flux at constant SiO2 content in the flux. These values vary widely, depending on other alloy concentrations. The data imply that manganese is almost always lost to the slag, and that the activity of manganese varies considerably with changes in the amount of titania in the flux. Manganese is very important to weld metal hardenability and must be closely controlled to obtain the optimum weld microstructure. Controlling the weld metal manganese concentration in the titania-containing flux systems would require strict compositional control of the welding flux to ensure the correct manganese concentration in the weld metal.

FIG. 3 PLOT OF CHANGES IN MANGANESE CONTENT IN THE WELD VERSUS THE VARIATION IN FLUX COMPOSITION AS A FUNCTION OF TWO-HEAT INPUTS. THE SIO2 CONTENT WAS MAINTAINED AT 40% THROUGHOUT THE FLUX-CORED ARC WELDING PROCESS. SOURCE: REF 21

In the same system, the delta titanium showed a constant increase with increasing titania content of the flux (Fig. 4). Because of the large amounts of titania in the flux, it is not surprising that the welds show positive compositional deviations for titanium. These results also suggest that control of the weld metal titanium content will require tight control of the limits of the TiO2 content of the flux.

FIG. 4 PLOT OF CHANGES IN TITANIUM CONTENT IN THE WELD VERSUS THE VARIATION IN FLUX COMPOSITION AS A FUNCTION OF TWO HEAT INPUTS. THE SIO2 CONTENT WAS MAINTAINED AT 40% THROUGHOUT THE FLUX-CORED ARC WELDING PROCESS. SOURCE: REF 21

If the delta quantity changes rapidly with flux composition, it may be difficult to maintain specific weld metal composition, microstructure, and properties with variation in flux composition. The magnitude of the delta quantity for specific elements is most often not as serious a concern as a rapid change in the delta quantity with variations in flux chemistry and welding parameters. The magnitude of the delta quantity can be adjusted by altering the alloy content of the welding wire or the amount of ferro-additions to the flux. Thus, a suitable combination of wire, flux, and welding parameters should achieve negligibly small delta quantities of major alloying elements. Shielding Gas. When the shielding gas that protects the weld pool comes from the flux, it is necessary to understand the

decomposition of specific flux components. Two common shielding gas atmospheres from flux dissociation are hydrogen and CO/CO2. The hydrogen gas can be produced by the decomposition of cellulose (wood flour or similar hydrocarbons). A CO/CO2 atmosphere results from the decomposition of carbonates such as limestone (CaCO3) or dolomite [CaMg(CO3)2]. The CO/CO2 atmosphere can be balanced to provide a reducing (and thus protective) atmosphere. At high temperatures, CO2 or CO will react with carbon to achieve equilibrium, which requires the presence of CO and CO2. It is the relative amounts of CO and CO2 that determine the reducing/oxidizing nature of the arc environment. The CO/CO2 ratio also determines the recovery of specific alloying elements. Typical plasma atmospheres for both hydrogentype and low-hydrogen-type electrodes are given in Table 2.

TABLE 2 GAS COMPOSITION OF WELDING ARC OBTAINED FROM SPECIFIC TYPES OF WELDING ELECTRODES

AWS TYPE COMPOSITION, WT% DESIGNATION H2O H2 CO2 CO 6010 CELLULOSIC 16 41 3 40

6015

BASIC

2

2

19

77

Source: Ref 15

In submerged arc welding, the covering flux also produces the protective shielding gas and plasma. Other gaseous phases, including flourine-bearing components, are also part of the plasma. The effect of arc environment (for example, amounts of CO and CO2) on weld metal chemistry control for hyperbaric welding with a basic electrode is shown in Fig. 5. The high-pressure welding allows evaluation of the CO reaction. Considering the CO reaction:

C + O = CO

(EQ 13)

the law of mass action gives:

(EQ 14) where [C] and [O] are the weld metal carbon and oxygen contents, respectively. At equilibrium the partial pressure of CO is directly related to the total pressure by Dalton's Law.

FIG. 5 EFFECT OF PRESSURE ON THE PRODUCT M = [%C][%O] FOR HYPERBARIC WELDING WITH A BASIC ELECTRODE. SOURCE: REF 12

In Fig. 5, two lines are indicated. The solid line plots the product of the weld metal oxygen and carbon as a function of total pressure when the analytical weld metal oxygen concentration, [%O]anal, was used. The broken line indicates a similar relationship; however, in this case, the weld metal oxygen concentration has been corrected for the displacement of oxygen because of the formation of manganese silicate inclusions during welding and its transport to the slag. The equilibrium oxygen content, [%O]eq, in the liquid steel at high temperatures is then given by the following expression:

(EQ 15) where ∆[%Si] and ∆[%Mn] represent the difference between the expected compositions of these elements (from knowledge of initial consumables and base plate compositions as well as dilution) and the actual compositions measured. The linearity of these lines in Fig. 5 is evidence of the strong influence of the CO reaction in arc welding. The evidence of water reaction control for some electrodes (for example, cellulosic electrodes) can be seen in Fig. 6. With an increase in the degree of oxidation of the deposited metal, the weld metal hydrogen content decreases (Ref 22). The relationship between weld metal oxygen and hydrogen contents as seen in Fig. 6 is consistent with the functional form expected from the law of mass action for a H2O reaction. It is apparent that with increasing weld metal oxygen content there is a significant reduction in weld metal hydrogen. The hydrogen content of the deposited metal can be reduced by increasing the CaCO3 content of the coating. With the ever-increasing requirements of weld-deposit properties and the increasing need for higher productivity, pyrochemistry will play an important role for flux formulation and weld-property prediction.

FIG. 6 PLOT OF WELD METAL OXYGEN CONTENT VERSUS WELD METAL HYDROGEN CONTENT WHEN WELDING WITH ELECTRODES THAT CONTAIN CHROMIUM AND NIOBIUM IN THEIR COATINGS. SOURCE: REF 22

Basicity Index As indicated in the section "Metal Transferability During Pyrochemical Reactions" in this article, the transfer of alloying elements during welding depends strongly on the physical and chemical properties of the flux. The ability to correlate flux properties with weld metal chemical compositions and properties is essential to understand the interactions between weld metal and flux. Because of the incomplete understanding of the thermodynamic properties of slags, the empirical concept of basicity has been applied to predict flux and weld properties. A basicity index, BI, for welding has been proposed:

(EQ 16)

where the chemical components are given in wt%. When the basicity index for a given flux is less than one, the flux is considered acidic. A flux with a BI between 1.0 and 1.2 can be classified as neutral. A flux with a BI greater than 1.2 is considered basic. In general, the higher the basicity, the cleaner the weld with respect to nonmetallic inclusions (that is, lower weld metal oxygen content) (Ref 23). Figure 7 illustrates the correlation between weld metal oxygen and the basicity index for some flux systems. The weld metal oxygen content drops significantly as the BI is increased to 1.2 and then remains relatively constant at about 250 ppm O. The correlation between weld metal oxygen, which is an indication of weld metal toughness, and basicity is acceptable for some welding flux systems, especially those that are primarily based on CaO, MgO, and SiO2. However, Eq 16 cannot be used to correlate the strength and toughness of welds made with high flux concentrations of Al2O3, TiO2, ZrO2, MnO, FeO, and CaF2. Although numerous basicity formulas have been considered, none has been flexible enough to deal with high amphoteric oxide contents. There is still some concern over utilizing this index (which was primarily conceived by steelmakers for evaluating sulfur refinement) for predicting weld metal oxygen, or for use as a general criterion for weld metal quality (Ref 7). Additionally, the BI does not consider the physical properties of the fluxes, nor does it explain the kinetics of the flux-metal reactions.

FIG. 7 EFFECT OF WELD METAL OXYGEN CONTENT ON FLUX BASICITY INDEX WHEN USING THE SUBMERGED ARC WELDING PROCESS. SOURCE: REF 23

Basicity index has been used by the welding industry as a measure of expected weld metal cleanliness and mechanical properties. Consequently, manufacturers of welding consumables have classified and advertised their fluxes with this index. It is believed that high basicity means high toughness, a quality of great interest to the engineer, while an acidic flux means excellent slag behavior, a characteristic of interest to the welder attempting to improve weld bead morphology and deposition rate. Table 3 classifies the various coating formulations for SMAW electrodes using the descriptors of cellulosic, basic, acid, and rutile (alternating current stabilizer and slag former).

TABLE 3 ELECTRODE COATING FORMULATIONS OF SELECTED SMAW ELECTRODES

Electrode Type

Coating formulation(a) AWS designation

Comments

Rutile

Cellulose

Quartz

Carbonates

Ferromanganese

Organics

Iron oremanganese ore

Calcium carbonate

Complex silicates

Fluorspar

Ferroalloys

Cellulosic 6010

2060

10-50

1530

0-15

5-10

...

...

...

...

...

...

Rutile

6012 6013

4060 2040

... ...

1525 1525

0-15 5-25

10-12 1.4-14

2-6 0-5

... ...

... ...

... ...

... ...

... ...

Acidore(a)

6020

...

...

X

X

X

...

X

...

X

...

...

Basic

7015

010

...

0-5

...

...

...

...

20-50

...

20-40

5-10

Source: Ref 24

(A)

X, DATA UNAVAILABLE.

Cellulose promotes gas shielding in the arc region. Hydrogen increases heat at weld. High hydrogen content (30-200 ppm). Deep penetration, fast cooling weld Slags mainly for slag shielding relatively high hydrogen content (15-30 ppm). High inclusion content in weld deposit Relatively high hydrogen content. High slag content in weld metal Relatively low hydrogen levels ( 10 ppm), hence commonly used in welding low-alloy construction steels. Electrodes should be kept dry. Low inclusion content in weld deposit

Pyrochemical Kinetics During Welding The ability of a flux to refine as well as protect the weld pool is related to the mass transport processes in the flux. The flux should melt approximately 200 °C (360 °F) below that of the alloy for proper flux coverage and for protection of the weld deposit. One of the most important physical properties of a flux is its slag viscosity, which not only governs the way the slag flows and covers the molten weld pool but also strongly affects the transport processes involved in pore removal, deoxidation, and retention of alloying additions. The chemical processing and refining by the flux to achieve a weld deposit with low concentrations of oxygen and sulfur and optimal concentration of hardenability agents (carbon, manganese, chromium, molybdenum, nickel, and so on) may not be achieved unless slag viscosity is also adequate. The viscosity is strongly temperature-dependent, so the use of various heat inputs during welding may require different flux compositions to produce the matching slag viscosity. For simple ionic melts, the viscosity has been shown to follow an Arrhenius-type temperature dependence:

(EQ 17) where ηis the viscosity, η0 is a system constant, Eη is the viscous activation energy, RT is the gas constant, and T is the temperature. For polymeric melts, Eq 17 often does not hold true. It has been shown that in certain cases a modified equation can be used:

(EQ 18)

where T0 is a constant for a given flux composition. The presence of entrapped particulates tends to increase the slag viscosity. Slag viscosity is also affected by composition. The compositional dependence is commonly reported by considering Eη to be a function of composition at constant temperature. The slag must be fluid enough so that it flows and covers the molten weld pool but must be viscous enough so that it does not run away from the molten metal and flow in front of the arc, leading to possible overlapping by the weld metal. (For overhead welding, surface tension becomes a primary factor because fluidity reduces coverage [opposite gravitational vector].) It has been reported that if the manganese silicate flux viscosity at 1450 °C (2640 °F) is above 0.7 Pa · s (7 P), a definite increase in weld surface pocking will occur. Pock marks have been associated with easily reducible oxides in the flux, which contribute oxygen to the weld pool. The weld pool reacts with carbon to form carbon monoxide, which cannot be transported through a high-viscosity flux and is trapped at the liquid-metal/flux interface. The result is a weld metal surface blemished by surface defects or pocks. Because viscosity is sensitive to temperature and thus heat input, pocking can be the evidence that a flux formulated for high-current welding is being used at too low a current or too great a travel speed. The viscosity of most welding fluxes at 1400 °C (2550 °F) is in the range of 0.2 to 0.7 Pa · s (2 to 7 P). Slag viscosity also affects the shape of the weld deposit and must be carefully controlled when covered electrodes are used out of position. The higher the slag viscosity, the greater the weld penetration in submerged arc welding. However, this benefit must be balanced, because if the viscosity is too high, the gaseous products cannot escape the weld pool, resulting in unacceptable porosity. This condition can be monitored by observing the density of pores trapped in the underside of the detached slag. Detached slags manifesting a honeycomb structure suggest a severe weld metal porosity problem. This condition usually means that a given flux has experienced an insufficient heat input for the effective transport of gas through the slag.


3. T. LAU, G.C. WEATHERLY, AND A. MCLEAN, THE SOURCES OF OXYGEN AND NITROGEN CONTAMINATION IN SUBMERGED ARC WELDING USING CAO-AL2O3 BASED FLUXES, WELD. J., VOL 64 (NO. 12), 1985, P 343S-347S

4. T.H. NORTH, H.B. BELL, A. NOWICKI, AND I. CRAIG, SLAG/METAL INTERACTION, OXYGEN, AND TOUGHNESS IN SUBMERGED ARC WELDING, WELD. J., VOL 57 (NO. 3), 1978, P 63S-75S 5. N. CHRISTENSEN AND J. CHIPMAN, SLAG-METAL INTERACTION IN ARC WELDING, WELD. RES. BULL., NO. 15, JAN 1953, P 1-14 6. R.H. FROST, D.L. OLSON, AND S. LIU, PYROCHEMICAL EVALUATION OF WELD METAL INCLUSION EVOLUTION, PROC. 3RD INT. CONF. TRENDS IN WELDING, ASM INTERNATIONAL, JUNE 1992 7. C.A. NATALIE, D.L. OLSON, AND M. BLANDER, PHYSICAL AND CHEMICAL BEHAVIOR OF WELDING FLUXES, ANN. REV. MATER. SCI., VOL 16, 1986, P 389-413 8. J.E. INDACOCHEA, M. BLANDER, N. CHRISTENSEN, AND D.L. OLSON, CHEMICAL REACTIONS WITH FEO-MNO-SIO2 FLUXES, METALL. TRANS. B, VOL 16, 1985, P 237-245 9. U. MITRA AND T.W. EAGAR, SLAG-METAL REACTIONS DURING WELDING, METALL. TRANS. B, VOL 22, 1991, P 65-100 10. C.S. CHAI AND T.W. EAGAR, SLAG-METAL EQUILIBRIUM DURING SUBMERGED ARC WELDING,METALL. TRANS. B, VOL 12, 1981, P 539-547 12. O. GRONG, D.L. OLSON, AND N. CHRISTENSEN, CARBON OXIDATION IN HYPERBARIC MMA WELDING, MET. CONSTRUCT., VOL 17. DEC 1985, P 810R-814R 14. T.W. EAGAR, SOURCES OF WELD METAL OXYGEN CONTAMINATION DURING SUBMERGED ARC WELDING, WELD. J.,VOL 57, 1978, P 76S-80S 15. N. CHRISTENSEN, WELDING METALLURGY, LECTURE NOTES, NTH, 1979 16. O. GRONG, T.A. SIEWERT, T.A. MARTINS, AND D.L. OLSON, A MODEL FOR THE SILICONMANGANESE DEOXIDATION OF STEEL WELD METAL, METALL. TRANS. A, VOL 17 (NO. 10), 1985, P 1797-1807 17. T.H. NORTH, THE DISTRIBUTION OF MANGANESE BETWEEN SLAG AND METAL DURING SUBMERGED ARC WELDING, WELD. RES. ABROAD, VOL 23 (NO. 1), 1977, P 2-40 18. T. BONISZEWSKI, BASIC FLUXES AND DEOXIDATION IN SUBMERGED ARC WELDING OF STEEL, METAL. CONSTR. BRIT. WELD. J., VOL 6, 1974, P 128 19. L. DAVIS, AN INTRODUCTION TO WELDING FLUXES FOR MILD AND LOW ALLOY STEELS, WELDING INSTITUTE, 1981 20. MAKING, SHAPING AND TREATING OF STEEL, UNITED STATES STEEL CORPORATION, 1984 21. P.S. DUNN, C.A. NATALIE, AND D.L. OLSON, SOL-GEL FLUXES FOR FLUX CORED WELDING CONSUMABLES, J. MATER. ENERGY SYSTEMS, VOL 8 (NO. 2), 1986, P 176-184 22. L.I. SOROKIN AND Z.A. SIDLIN, THE EFFECT OF ALLOYING ELEMENTS AND OF MARBLE IN AN ELECTRODE COATING ON THE SUSCEPTIBILITY OF A DEPOSITED NICKEL CHROME METAL TO PORE FORMATION, SVAR. PROIZVOD., NO. 11, 1974, P 7-9 23. S.S. TULIANI, T. BONISZEWSKI, AND N.F. EATON, NOTCH TOUGHNESS OF COMMERCIAL SUBMERGED ARC WELD METAL, WELD. MET. FABR., VOL 37 (NO. 8), 1969, P 27 24. T.G.F. GRAY, J. SPENCE, AND T.H. NORTH, RATIONAL WELDING DESIGN, BUTTERWORTHS, 1975 Nature and Behavior of Fluxes Used for Welding D.L. Olson, S. Liu, R.H. Frost, G.R. Edwards, and D.A. Fleming, Colorado School of Mines

Alloy Modification Other additions to the welding consumables are required to make alloy additions to the weld pool, usually in the form of powder metal or ferro-additions (Ref 1). Often the composition of the wire that makes up the rod for the shielded metal arc electrode from a specific manufacturer is the same, regardless of the alloy to be welded. Alloying is achieved by

powder metal additions to the flux coating. Manganese, silicon, chromium, niobium, and other alloying additions are adjusted in the weld pool by ferroalloy powder additions. Specially prepared alloy additions of Fe-50Si, Fe-80Mn, Fe60Mn-30Si, and others are used. One concern in formulating electrodes using ferro-additions is the alloying element recovery (that is, the amount of the element that is transferred across the arc and into the weld deposit). Values for the recovery of typical elements in steel welding are given in Table 4. The metal losses are either to the slag or to the fume.

TABLE 4 RECOVERY OF ELEMENTS FROM SELECTED ELECTRODE COVERINGS

ALLOY ELEMENT FORM OF MATERIAL IN ELECTRODE COVERING ALUMINUM FERROALUMINUM BORON FERROBORON CARBON GRAPHITE CHROMIUM FERROCHROMIUM NIOBIUM FERROCOLUMBIUM COPPER COPPER METAL MANGANESE FERROMANGANESE MOLYBDENUM FERROMOLYBDENUM NICKEL ELECTROLYTIC NICKEL SILICON FERROSILICON TITANIUM FERROTITANIUM VANADIUM FERROVANADIUM

APPROXIMATE RECOVERY OF ELEMENT, WT% 20 2 75 95 70 100 75 97 100 45 5 80

Source: Ref 1 Slipping and Binding Agents. In the case of SMAW electrodes, slipping agents are also added to the green flux formulation to improve the extrudability of the flux onto the rod (Ref 1). Glycerin, china clay, kaolin, talc, bentonite, and mica have all been used as slipping agents. Binding agents can be classified into two types. The first type comprises binders that bond the flux components to the rod without introducing a hydrogen source. These low-hydrogen binders include sodium silicate and potassium silicate. The second type of binding agent does function as a hydrogen source. The binders used for high-hydrogen electrodes can be organic in nature and include gum arabic, sugar, dextrine, and other specialized synthetic organic binders. Slag Formation. Slag, a mixture of glass and crystalline structure, must solidify on the already solidified weld deposit to

protect the surface from oxidation during cooling. Specific physical properties are required of the slag. It must melt below the melting temperature of steel (~ 1450 °C, or 2640 °F), must have a density significantly less than steel to reduce slag entrapment in the weld deposit, must possess the proper viscosity in the temperature range of 1450 to 1550 °C (2640 to 2820 °F), and must easily detach from the weld deposit after welding. Silicates, aluminates, and titanates are all primary slag formers. The high-valence cations of these compounds produce a bonding network that can promote glass formation. Most electrodes produce silicate or titanate slag. Silicates of such elements as manganese produce very workable glass deposits, but also produce weld deposits that are relatively high in weld metal oxygen content. These fluxes are said to range from acidic to neutral. The titanate and aluminate fluxes produce more rigid and stable oxide covers. The result is a lower concentration of weld metal oxygen, but a higherviscosity slag. Aluminates and calcium-bearing compounds are common additions to basic fluxes for the submerged arc welding of high-toughness linepipe steels. Minerals used for slag formation include: • • • • •

RUTILE POTASSIUM TITANATE ILMENITE ALUMINA SILICA FLOUR

• • • •

IRON POWDER FLUORSPAR FELDSPAR MANGANESE DIOXIDE

Asbestos (up to 50%) was used as a slag former, but has been phased out by welding consumable manufacturers. Slag detachability is a serious productivity concern for steel fabricators. The relative ease of slag detachability influences

the economic advantages of flux-related welding processes. Residual slag on the weld deposit promotes slag stringers in multipass weldments, limits the effective use of narrow-gap flux-related welding processes, and may reduce the corrosion resistance of the weldment. Welding flux formulators have modified flux compositions to alleviate or reduce this hindrance (Ref 25). Poor slag removal has been reported to occur when the flux contains fluorite. Slags containing spinels generally have been found to attach tenaciously to the weld deposit. Slags with cordierite and (Cr,Mn,Mg)O · (Cr,Mn,Al)2O3 type spinel phases have been reported to be difficult to remove from stainless steel weldments. It has also been reported to be difficult to remove from stainless steel weldments. It has also been reported that if (CaO)2SiO2, Cr2TiO5, and FeTiO5 are present, the slag readily detaches from the weld deposit (Ref 25). The compositional range for acceptable slag detachability has been reported for the CaO · CaF2 · SiO2 system and the CaO · TiO2 · SiO2 system. Easier slag removal has also been related to deoxidation with aluminum instead of titanium. Increasing the Al2O3 content in the flux has demonstrated improved detachability (Ref 25). Differences in the coefficient of thermal expansion between the slag and the weld are also important.


1. G.E. LINNERT, CHAPTER 8, WELDING METALLURGY, VOL 1, AWS, 1965, P 367-396 25. D.L. OLSON, G.R. EDWARDS, AND S.K. MARYA, THE PHYSICAL AND CHEMICAL BEHAVIOR ASSOCIATED WITH SLAG DETACHABILITY DURING WELDING, FERROUS ALLOY WELDMENTS, VOL 67-70, TRANS TECH, 1992, P 253-268

Nature and Behavior of Fluxes Used for Welding D.L. Olson, S. Liu, R.H. Frost, G.R. Edwards, and D.A. Fleming, Colorado School of Mines

Types of Fluxes Producing a weld deposit with high mechanical integrity and with high productivity requires many physical and chemical occurrences in the welding arc. These occurrences are achieved in flux-related arc welding processes by the careful formulation of the welding flux. The welding flux must: • • • • • • • • •

STABILIZE ARC AND CONTROL ARC RESISTIVITY PROVIDE SLAG WITH PROPER MELTING TEMPERATURE PROVIDE LOW-DENSITY SLAG PERMIT USE OF DIFFERENT TYPES OF CURRENT AND POLARITY ADD ALLOYING ELEMENTS REFINE THE WELD POOL (DEOXIDATION AND DESULFURIZATION) PROVIDE PROPER VISCOSITY FOR OUT-OF-POSITION WELDING PROMOTE SLAG DETACHABILITY PRODUCE SMOOTH WELD CONTOUR

•

REDUCE SPATTER AND FUME

It is apparent, after considering the large number of necessary requirements, that a welding flux must be carefully and deliberately formulated to achieve an optimized performance. A number of different arc welding processes depend on welding fluxes. Each of these processes requires a different formulation. SMAW Fluxes. The typical constituents and their functions in electrode coatings for the SMAW process are given in

Table 5. It should be noted that the flux ingredients are based on additions of refined minerals of the earth. Natural minerals offer an economical method of keeping welding consumables at a reasonable cost. Table 6 presents the elemental content of these minerals. Table 7 gives the typical chemical compositions for flux coatings for three different SMAW electrodes.

TABLE 5 TYPICAL FUNCTIONS AND COMPOSITIONS OF CONSTITUENTS FOR SELECTED MILD STEEL SMAW ELECTRODE COATINGS

COATING CONSTITUENT

FUNCTION OF CONSTITUENT PRIMARY SECONDARY

CELLULOSE

SHIELDING GAS SHIELDING GAS SLAG FORMER SHIELDING GAS SLAG FORMER ARC STABILIZER SLAG FORMER EXTRUSION EXTRUSION

CALCIUM CARBONATE FLUORSPAR DOLOMITE TITANIUM DIOXIDE (RUTILE) POTASSIUM TITANATE FELDSPAR MICA CLAY SILICA ASBESTOS MANGANESE OXIDE IRON OXIDE IRON POWDER FERROSILICON FERROMANGANESE SODIUM SILICATE

SLAG FORMER SLAG FORMER SLAG FORMER SLAG FORMER DEPOSITION RATE DEOXIDIZER ALLOYING BINDER

... FLUXING AGENT FLUXING AGENT FLUXING AGENT ARC STABILIZER SLAG FORMER STABILIZER

COMPOSITION RANGE OF COATING ON ELECTRODE E6010, E6013 E7018 E7024 E7028 E6011 25-40 2-12 ... 1-5 ... ...

0-5

15-30

0-5

0-5

...

...

15-30

...

5-10

...

...

...

...

5-10

10-20

30-55

0-5

20-35

10-20

(A)

(A)

0-5

...

0-5

...

0-20

0-5

...

0-5

STABILIZER SLAG FORMER ...

... ...

0-15 0-10

... ...

0-15 ...

... ...

...

...

...

...

...

EXTRUSION

10-20

...

...

...

...

ALLOYING

...

...

...

...

...

...

...

...

...

...

CONTACT WELDING ... DEOXIDIZER FLUXING

...

...

25-40

40-55

40-55

... 5-10 20-30

... 5-10 5-10

5-10 2-6 0-5

0-5 5-10 0-10

2-6 2-6 0-5

POTASSIUM SILICATE

ARC STABILIZER

AGENT BINDER

(B)

5-15(B) 5-10

0-10

0-5

(A) REPLACES TITANIUM DIOXIDE (RUTILE) TO PERMIT USE WITH ALTERNATING CURRENT. (B) REPLACES SODIUM SILICATE TO PERMIT USE WITH ALTERNATING CURRENT TABLE 6 TYPICAL COMPOSITION OF COMMON MINERALS USED IN SMAW ELECTRODE COATINGS

MINERAL ILMENITE TALC BENTONITE SILICA, QUARTZ CELLULOSE ALUMINA MUSCOVITE, MICA(A) ACTINOLITE MAGNETITE HEMATITE RUTILE, TITANIA DOLOMITE FLUORSPAR, FLUORITE CRYOLITE LIME LIMESTONE, CALCITE, MARBLE ZIRCONIA FELDSPAR(A) CLAY(A) SODIUM SILICATE POTASSIUM SILICATE CHROMIC OXIDE

CHEMICAL COMPOSITION FEO · TIO2 3MGO · 4SIO2 · H2O COMPLEX AL, MG, CA, FE HYDROXIDES SIO2 (C6H10O5)X AL2O3 K2O · 3AL2 · 6SIO2 · 2H2O CAO · MGO · 2FEO · 4SIO2 FE3O4 FE2O3 TIO2 MGO · CAO · (CO2)2 CAF2 NA3ALF4 CAO CACO3 ZRO2 K2O · AL2O3 · 6SIO2 AL2O3 · 2SIO2 · 2H2O SIO2/NA2O (RATIO-3.22) SIO2 K2O (RATIO-2.11) CR2O3

Source: Ref 1

(A) ALTHOUGH THESE SUBSTANCES CAN HAVE SEVERAL CHEMICAL COMPOSITIONS, ONLY TYPICAL COMPOSITION IS GIVEN.

TABLE 7 CHEMICAL COMPOSITION OF COVERINGS USED IN ELECTRODES FOR SMAW WELDING OF MILD STEELS AND LOW-ALLOY STEELS

COMPOSITION, WT%(A)

ELECTRODE DESCRIPTION

AWS DESIGNATION

HIGH-CELLULOSE, GAS E6010 SHIELDED HIGH-TITANIUM GAS-SLAG E6012 SHIELD LOW-HYDROGEN IRON E7018 POWDER Source: Ref 1

(A)

AFTER BAKING.

CAO

TIO2

...

10.1 . . .

...

46.0 . . .

14.4 . . .

CAF2

SIO2

AL2O3

MGO

NA3ALF3

N3O

FEO

47.0 . . .

3.2

...

5.1 1.3

23.6 5.0

2.0

...

11.0 20.5 2.0

1.0

5.0

SI

MN

FE

CO + CO2

VOLATILE MATTER

MOISTURE

1.5 2.8 . . .

...

25.0

4.0

2.4 7.0

1.5 2.5 . . .

...

5.0

2.0

1.2 . . .

2.5 1.8 28.5 12.0

...

0.1

Flux-cored arc welding (FCAW) uses a hollow wire filled with flux reagents and ferro-additions. The two types of

flux-cored electrodes are carbon-dioxide-shielded flux-cored electrodes and self-shielded flux-cored electrodes (Ref 26). Table 8 gives typical compositions for the three types of carbon-dioxide-shielded cored electrodes. Table 9 gives the typical compositions for the four types of self-shielded flux-cored electrodes.

TABLE 8 TYPICAL FLUX COMPOSITIONS OF AVAILABLE CO2 SHIELDED FCAW ELECTRODES

ELECTRODE DESCRIPTION FLUX TYPE NUMB TYPE ER 1 TITANIA (NONBAS IC) 2 LIMETITANIA (BASIC OR NEUTRA L) 3 LIME (BASIC)

COMPOSITION, WT% AWS DESIGNAT ION

SIO2

E70T-1, E70T-2

AL2 O3

TIO 2

ZR O2

CA O

21. 0

2.1

40. 5

. . . 0. 7

E70T-1

17. 8

4.3

9.8

E70T-1, E70T-5

7.5

0.5

...

NA2 O

K2 O

CO2 (AS CARBONA TE)

C

1.6 1. 4

0.5

0. 6

20. 1

15. 8

...

6.2 9. 7

1.9 1. 5

...

0. 3

24. 7

13. 0

18. 0

. . . 3. 2

. . . 0. 5

2.5

1. 1

55. 0

7.2

20. 5

FE

MN

CAF 2

Source: Ref 26

TABLE 9 TYPICAL FLUX COMPOSITIONS OF AVAILABLE SELF-SHIELDED FCAW ELECTRODES

ELECTRODE DESCRIPTION TYPE FLUX NUMB TYPE ER 1 FLUORSP ARALUMIN UM 2 FLUORSP ARTITANIA 3 FLUORSP AR-LIMETITANIA 4 FLUORSP AR-LIME

COMPOSITION, WT% AWS DESIGNAT ION

SI O2

AL

E70T-4, E70T-7, E60T-8

0. 5

15. . . . 4

. . . . . 12. 0. . 6 4

E70T-3

3. 6

1.9 . . .

20. . . 4.5 0. 6 . 6

E70T-6

4. 2

1.4 . . .

E70T-5

6. 9

. . . 0.6

AL2 O3

CO2 (AS CARBON ATE)

FE

0.2 1. 2

0.4

0.4 3. 0

. . 63. . 5

0.1 0. 6

0.6

50. 4. 0 5

. . 22. . 0

14. 4. 7 0

2.2 . . . . . 0. . 6

2.1

50. 2. 5 0

2. 4

1.2 3. 2

. . . . . 0.6 0. . 3

1.3

58. 7. 0 9

. . 22. . 0

TIO 2

CA O

MG O

K2 O

NA2 O

C

M N

NI

CA F2

15. 3

Source: Ref 26 Submerged Arc Welding (SAW). In the SAW process, the flux drops from a hopper onto the work such that the welding arc is submerged beneath the granular flux, producing an arc cavity that contains metal vapors, silicon monoxide, manganese oxide, gaseous fluorides, and other higher-vapor-pressure components of the flux. This arc welding process has been recognized as one that produces very little, if any, fume. Submerged arc welding is limited mainly to the flat or

horizontal position, and requires significant setup time. It is very successful in manufacturing numerous similar parts, such as producing welded steel pipe. Submerged arc fluxes are made in three different forms: • • •

BONDED FLUXES MIX NONMETALLIC AND FERRO-ADDITIONS WITH A LOWTEMPERATURE BINDER INTO MIXTURES OF SMALL PARTICLES. AGGLOMERATED FLUXES ARE SIMILAR TO THE BONDED FLUXES EXCEPT THAT A CERAMIC GLASS BINDER, CURED AT HIGH TEMPERATURE, IS USED. FUSED FLUXES ARE MADE BY POURING A HOMOGENEOUS GLASS MIXTURE OF THE PROPER FLUX COMPOSITION INTO WATER, RESULTING IN A FRIT.

Generally speaking, changing from bonded to agglomerated to fused fluxes improves control of the weld metal composition, especially with respect to such impurities as hydrogen and oxygen. There are seven types of submerged arc fluxes (Table 10). Flux classification according to basicity is the result of observed correlations between weld metal oxygen concentration and flux composition. Table 11 gives some typical compositions for SAW fluxes.

TABLE 10 FLUXES USED FOR SAW APPLICATIONS

FLUX TYPE MANGANESE SILICATE

CONSTITUENTS MNO + SIO2 > 50%

BASICITY FLUX FORM ACID FUSED

ADVANTAGES MODERATE STRENGTH; TOLERANT TO RUST; FAST WELDING SPEEDS; HIGH HEAT INPUT; GOOD STORAGE

CALCIUMHIGH SILICA

CAO + MGO + SIO2 > 60%

ACID

AGGLOMERATED, FUSED

HIGH WELDING CURRENT; TOLERANT TO RUST

CALCIUM SILICATENEUTRAL


NEUTRAL

AGGLOMERATED, FUSED

CALCIUM SILICATELOW SILICA


BASIC

AGGLOMERATED, FUSED

ALUMINATE

AL2O3 + CAO +

BASIC

AGGLOMERATED

MODERATE STRENGTH AND TOUGHNESS; ALL CURRENT TYPES; TOLERANT TO RUST; SINGLE- OR MULTIPLE-PASS WELD GOOD TOUGHNESS WITH MEDIUM STRENGTH; FAST WELDING SPEEDS; LESS CHANGE IN COMPOSITION AND LOWER OXYGEN GOOD STRENGTH

LIMITATIONS LIMITED USE FOR MULTIPASS WELDING; USE WHERE NO TOUGHNESS REQUIREMENT; HIGH WELD METAL OXYGEN; INCREASE IN SILICON ON WELDING; LOW IN CARBON POOR WELD TOUGHNESS; USE WHERE NO TOUGHNESS REQUIREMENT; HIGH WELD METAL OXYGEN ...

COMMENTS ASSOCIATED MANGANESE GAIN; MAXIMUM CURRENT, 1100 A; HIGHER WELDING SPEEDS

NOT TOLERANT TO RUST; NOT USED FOR MULTIWIRE WELDING

...

NOT TOLERANT TO

USUALLY

DIFFER IN SILICON GAIN; SOME CAPABLE OF 2500 A; WIRES WITH HIGH MANGANESE ...

BASIC

MGO > 45%; AL2O3 > 20%

ALUMINA

BAUXITE BASE

NEUTRAL

AGGLOMERATED, FUSED

BASIC FLUORIDE

CAO + MGO + MNO + CAF2 > 50%; SIO2 22%; CAF2 15%

BASIC

AGGLOMERATED, FUSED

AND TOUGHNESS IN MULTIPASS WELDS; NO CHANGE IN CARBON; LOSS OF SULFUR AND SILICON LESS CHANGE IN WELD COMPOSITION AND LOWER OXYGEN THAN FOR ACID TYPE; MODERATE TO FAST WELDING SPEEDS VERY LOW OXYGEN; MODERATE TO GOOD LOWTEMPERATURE TOUGHNESS

RUST, LIMITED TO DC ELECTRODE POSITIVE

...

MAY PRESENT PROBLEMS OF SLAG DETACHABILITY; MAY PRESENT PROBLEM OF MOISTURE PICKUP

MANGANESE GAIN; MAXIMUM CURRENT 1200 A; GOOD MECHANICAL PROPERTIES ...

CAN BE USED WITH ALL WIRES, PREFERABLE DC WELDING; VERY GOOD WELD PROPERTIES

TABLE 11 TYPICAL COMPOSITIONS OF SMAW FLUXES

FLUX COMPOSITION, WT% AL2O3 SIO2 TIO2 MGO A 49.9 13.7 10.1 2.9 B 24.9 18.4 0.2 28.9 C 19.3 16.3 0.8 27.2 D 18.1 13.2 0.5 28.2 E 17.0 12.2 0.7 36.8

CAF2 5.7 24.2 23.6 31.8 29.2

CAO ... ... 9.8 4.5 0.7

MNO 15.1 1.8 0.08 0.1 8.9

NA2O 1.6 2.1 0.9 0.9 1.6

K2O 0.2 0.07 1.1 0.9 0.1

BASICITY INDEX (BI) 0.4 1.8 2.4 3.0 3.5


1. G.E. LINNERT, CHAPTER 8, WELDING METALLURGY, VOL 1, AWS, 1965, P 367-396 26. FUMES AND GASES IN THE WELDING ENVIRONMENT, AWS, 1979 Nature and Behavior of Fluxes Used for Welding D.L. Olson, S. Liu, R.H. Frost, G.R. Edwards, and D.A. Fleming, Colorado School of Mines

References

1. G.E. LINNERT, CHAPTER 8, WELDING METALLURGY, VOL 1, AWS, 1965, P 367-396 2. C.E. JACKSON, FLUXES AND SLAGS IN WELDING, WELD. RES. BULL., NO. 190, 1973 3. T. LAU, G.C. WEATHERLY, AND A. MCLEAN, THE SOURCES OF OXYGEN AND NITROGEN CONTAMINATION IN SUBMERGED ARC WELDING USING CAO-AL2O3 BASED FLUXES, WELD. J., VOL 64 (NO. 12), 1985, P 343S-347S 4. T.H. NORTH, H.B. BELL, A. NOWICKI, AND I. CRAIG, SLAG/METAL INTERACTION, OXYGEN, AND TOUGHNESS IN SUBMERGED ARC WELDING, WELD. J., VOL 57 (NO. 3), 1978, P 63S-75S 5. N. CHRISTENSEN AND J. CHIPMAN, SLAG-METAL INTERACTION IN ARC WELDING, WELD. RES. BULL., NO. 15, JAN 1953, P 1-14 6. R.H. FROST, D.L. OLSON, AND S. LIU, PYROCHEMICAL EVALUATION OF WELD METAL INCLUSION EVOLUTION, PROC. 3RD INT. CONF. TRENDS IN WELDING, ASM INTERNATIONAL, JUNE 1992 7. C.A. NATALIE, D.L. OLSON, AND M. BLANDER, PHYSICAL AND CHEMICAL BEHAVIOR OF WELDING FLUXES, ANN. REV. MATER. SCI., VOL 16, 1986, P 389-413 8. J.E. INDACOCHEA, M. BLANDER, N. CHRISTENSEN, AND D.L. OLSON, CHEMICAL REACTIONS WITH FEO-MNO-SIO2 FLUXES, METALL. TRANS. B, VOL 16, 1985, P 237-245 9. U. MITRA AND T.W. EAGAR, SLAG-METAL REACTIONS DURING WELDING, METALL. TRANS. B, VOL 22, 1991, P 65-100 10. C.S. CHAI AND T.W. EAGAR, SLAG-METAL EQUILIBRIUM DURING SUBMERGED ARC WELDING,METALL. TRANS. B, VOL 12, 1981, P 539-547 11. N. CHRISTENSEN, METALLURGICAL ASPECTS OF ARC WELDING, WELD. J., VOL 27, 1949, P 373S-380S 12. O. GRONG, D.L. OLSON, AND N. CHRISTENSEN, CARBON OXIDATION IN HYPERBARIC MMA WELDING, MET. CONSTRUCT., VOL 17. DEC 1985, P 810R-814R 13. U. MITRA AND T.W. EAGAR, SLAG-METAL REACTIONS DURING SUBMERGED ARC WELDING OF ALLOY STEELS, METALL. TRANS. A, VOL 15, 1984, P 217-227

14. T.W. EAGAR, SOURCES OF WELD METAL OXYGEN CONTAMINATION DURING SUBMERGED ARC WELDING, WELD. J.,VOL 57, 1978, P 76S-80S 15. N. CHRISTENSEN, WELDING METALLURGY, LECTURE NOTES, NTH, 1979 16. O. GRONG, T.A. SIEWERT, T.A. MARTINS, AND D.L. OLSON, A MODEL FOR THE SILICONMANGANESE DEOXIDATION OF STEEL WELD METAL, METALL. TRANS. A, VOL 17 (NO. 10), 1985, P 1797-1807 17. T.H. NORTH, THE DISTRIBUTION OF MANGANESE BETWEEN SLAG AND METAL DURING SUBMERGED ARC WELDING, WELD. RES. ABROAD, VOL 23 (NO. 1), 1977, P 2-40 18. T. BONISZEWSKI, BASIC FLUXES AND DEOXIDATION IN SUBMERGED ARC WELDING OF STEEL, METAL. CONSTR. BRIT. WELD. J., VOL 6, 1974, P 128 19. L. DAVIS, AN INTRODUCTION TO WELDING FLUXES FOR MILD AND LOW ALLOY STEELS, WELDING INSTITUTE, 1981 20. MAKING, SHAPING AND TREATING OF STEEL, UNITED STATES STEEL CORPORATION, 1984 21. P.S. DUNN, C.A. NATALIE, AND D.L. OLSON, SOL-GEL FLUXES FOR FLUX CORED WELDING CONSUMABLES, J. MATER. ENERGY SYSTEMS, VOL 8 (NO. 2), 1986, P 176-184 22. L.I. SOROKIN AND Z.A. SIDLIN, THE EFFECT OF ALLOYING ELEMENTS AND OF MARBLE IN AN ELECTRODE COATING ON THE SUSCEPTIBILITY OF A DEPOSITED NICKEL CHROME METAL TO PORE FORMATION, SVAR. PROIZVOD., NO. 11, 1974, P 7-9 23. S.S. TULIANI, T. BONISZEWSKI, AND N.F. EATON, NOTCH TOUGHNESS OF COMMERCIAL SUBMERGED ARC WELD METAL, WELD. MET. FABR., VOL 37 (NO. 8), 1969, P 27 24. T.G.F. GRAY, J. SPENCE, AND T.H. NORTH, RATIONAL WELDING DESIGN, BUTTERWORTHS, 1975 25. D.L. OLSON, G.R. EDWARDS, AND S.K. MARYA, THE PHYSICAL AND CHEMICAL BEHAVIOR ASSOCIATED WITH SLAG DETACHABILITY DURING WELDING, FERROUS ALLOY WELDMENTS, VOL 67-70, TRANS TECH, 1992, P 253-268 26. FUMES AND GASES IN THE WELDING ENVIRONMENT, AWS, 1979

Shielding Gases for Welding Kevin A. Lyttle, Praxair, Inc.

Introduction THE SHIELDING GAS used in a welding process has a significant influence on the overall performance of the welding system. Its primary function is to protect the molten metal from atmospheric nitrogen and oxygen as the weld pool is being formed. The shielding gas also promotes a stable arc and uniform metal transfer. In gas-metal arc welding (GMAW) and flux-cored arc welding (FCAW), the gas used has a substantial influence on the form of metal transfer during welding. This, in turn, affects the efficiency, quality, and overall operator acceptance of the welding operation. The shielding gas interacts with the base material and with the filler material, if any, to produce the basic strength, toughness, and corrosion resistance of the weld. It can also affect the weld bead shape and the penetration pattern. Understanding the basic properties of a shielding gas will aid in the selection of the right shielding gas or gases for a welding application. Use of the best gas blend will improve the quality and may reduce the overall cost of the welding operation as well. Shielding Gases for Welding Kevin A. Lyttle, Praxair, Inc.

Basic Properties of a Shielding Gas The "controlled electrical discharge" known as the welding arc is formed and sustained by the establishment of a conductive medium called the arc plasma. This plasma consists of ionized gas, molten metals, slags, vapors, and gaseous atoms and molecules. The formation and structure of the arc plasma is dependent on the properties of the shielding gases used for welding. Table 1 lists the basic properties of gases used for welding (Ref 1).

TABLE 1 PROPERTIES OF SHIELDING GASES USED FOR WELDING

GAS

ARGON CARBON DIOXIDE HELIUM HYDROGEN NITROGEN OXYGEN

CHEMICAL MOLECULAR SPECIFIC DENSITY (A) SYMBOL WEIGHT GRAVITY g/ft3 g/l AR 39.95 1.38 0.1114 1.784 CO2 44.01 1.53 0.1235 1.978

IONIZATION POTENTIAL aj(b) eV 2.52 15.7 2.26 14.4

HE H2 N2 O2

3.92 2.16 2.32 2.11

4.00 2.016 28.01 32.00

0.1368 0.0695 0.967 1.105

0.0111 0.0056 0.782 0.0892

0.178 0.090 12.5 1.43

24.5 13.5 14.5 13.2

Source: Ref 1

(A) (B)

AT 100 KPA (1 ATM) AND 0 °C (32 °F); AIR = 1. 10-18 J.

The ionization potential is the energy, expressed in electron volts, necessary to remove an electron from a gas atom--

making it an ion, or an electrically charged gas atom. All other factors held constant, the value of the ionization potential decreases as the molecular weight of the gas increases. Arc starting and arc stability are greatly influenced by the ionization potentials of the component shielding gases used in welding process. A gas with a low ionization potential,

such as argon, can atoms into ions easily. Helium, with its significantly higher ionization potential, produces a harder to start, less stable arc. Although other factors are involved in sustaining the plasma, the respective energy levels required to ionize these gases must be maintained; as a consequence, the arc voltage is directly influenced. For equivalent arc lengths and welding currents, the voltage obtained with helium is appreciably higher than is with argon. This translates into more available heat input to the base material with helium than with argon. The thermal conductivity of a gas is a measure of how well it is able to conduct heat. It influences the radial heat loss

from the center to the periphery of the arc column as well as heat transfer between the plasma and the liquid metal. Argon, which has a low thermal conductivity, produces an arc that has two zones: a narrow hot core and a considerably cooler outer zone. The penetration profile of the weld fusion area then exhibits a narrow "finger" at the root and a wider top. A gas with a high thermal conductivity conducts heat outward from the core; this results in a wider, hotter arc core. This type of heat distribution occurs with helium, argon-hydrogen, and argon-carbon dioxide blends; it gives a more even distribution of heat to the work surface and produces a wider fusion area. Dissociation and Recombination. Shielding gases such as carbon dioxide, hydrogen, and oxygen are multiatom

molecules. When heated to high temperatures within the arc plasma, these gases break down, or dissociate, into their component atoms. They are then at least partially ionized, producing free electrons and current flow. As the dissociated gas comes into contact with the relatively cool work surface, the atoms recombine and release heat at that point. This heat of recombination causes multiatomic gases to behave as if they have a higher thermal conductivity, similar to that of helium. Dissociation and recombination do not occur with gases, such as argon, that consist of a single atom. Thus, at the same arc temperature, the heat generated at the work surface can be considerably greater with gases such as carbon dioxide and hydrogen. Reactivity/Oxidation Potential. The oxidizing nature of the shielding gas affects both welding performance and the

properties of the resultant weld deposit. Argon and helium are completely nonreactive, or inert, and thus have no direct chemical affect on the weld metal. Oxidizing or active gases, such as CO2 and oxygen, will react with elements in the filler metal or baseplate and will form a slag on the surface of the weld deposit. The loss of elements, such as manganese and silicon, from steel can affect the quality and cost of the weldment produced. Both weld strength and toughness generally decline as the oxidizing nature of the shielding gas increases. Additions of reactive gases such as oxygen or carbon dioxide enhance the stability of the arc and affect the type of metal transfer obtained. Metal droplet size is decreased, and the number of droplets transferred per unit time increases as the level of oxygen in the shielding gas increases. Oxygen reduces the molten weld bead surface tension, promoting better bead wetting and higher welding travel speeds. Small additions of CO2 work in a similar manner. The surface tension between the molten metal and its surrounding atmosphere has a pronounced influence on bead

shape. If the surface energy is high, a convex, irregular bead will result. Low values promote flatter beads with minimal susceptibility for undercutting. Pure argon is generally associated with high interfacial energy, producing a sluggish weld puddle and high, crowned bead. The addition of a small amount of a reactive gas, such as oxygen, lowers this surface tension and promote fluidity and better wetting of the base material; it does this without creating excessive oxidation of the weld metal. Gas Purity. Some metals, such as carbon steel and copper, have a relatively high tolerance for contaminants in the

shielding gas; others, such as aluminum and magnesium, are fairly sensitive to particular contaminants. Still others, such as titanium and zirconium, have an extremely low tolerance for any foreign constituent in the shielding gas. Depending on the metal being welded and the welding process used, very small quantities of gas impurities can significantly affect welding speed, weld surface appearance, weld bead solidification, and porosity levels. The effects of any given impurity are wide ranging, but weld quality and eventual fitness for purpose are major areas of concern. There is always a possibility that the gas, as delivered, is contaminated; however, it is far more likely that impurities will enters somewhere between the supply and the end-use points. For this reason, property designed piping systems and highquality hose are recommended for use with welding shielding gases. Typical industry minimum purity levels for welding gases are listed in Table 2 (Ref 2).

TABLE 2 TYPICAL GASES PURITY AND MOISTURE CONTENT OF SHIELDING GAS

ARGON CARBON DIOXIDE HELIUM HYDROGEN NITROGEN OXYGEN

PRODUCT STATE

GAS LIQUID GAS LIQUID GAS LIQUID GAS LIQUID GAS LIQUID INDUSTRIAL LIQUID

MINIMUM PURITY, %

99.995 99.997 99.5 99.8 99.95 99.995 99.95 99.995 99.7 99.997 99.5 99.5

MAXIMUM MOISTURE, PPM(A)

10 6 19 50 32 3 8 5 32 5 50 6

APPROXIMATE DEW POINT AT MAXIMUM MOISTURE CONTENT

°C -60 -64 -51 -58 -51 -69 -63 -65 -51 -65 -48 -64

°F -77 -83 -60 -73 -61 -92 -80 -86 -61 -86 -54 -83

Source: Ref 2

(A) MOISTURE SPECIFICATIONS ARE MEASURED AT FULL CYLINDER PRESSURE, THE PRESSURE AT WHICH THE CYLINDER IS ANALYZED. Gas density is the weight of the gas per unit volume. Density is one of the chief factors that influence shielding gas effectiveness. Basically, gases heavier than air, such as argon and carbon dioxide, require lower flow rates in use than do the lighter gases, such as helium, to ensure adequate protection of the weld puddle.


1. N.E. LARSON AND W.F. MEREDITH, SHIELDING GAS SELECTION MANUAL, UNION CARBIDE INDUSTRIAL GASES TECHNOLOGY CORP., 1990, P 10 2. N.E. LARSON AND W.F. MEREDITH, SHIELDING GAS SELECTION MANUAL, UNION CARBIDE INDUSTRIAL GASES TECHNOLOGY CORP., 1990, P 11

Shielding Gases for Welding Kevin A. Lyttle, Praxair, Inc.

Characteristics of the Components of a Shielding Gas Blend To obtain a shielding gas that is suited to a specific application, a mix of gases is generally needed. Each basic gas contributes certain characteristics to the performance of the overall mix. Some gas blends have relatively specific areas of application and limited operating ranges; others can be used on many materials under a variety of welding conditions. Each component of the blend brings with it properties that are supplemented by the others to produce an enhanced level of performance.

Argon is inert or unreactive with respect to the materials present in the welding electrode. With its low ionization

potential, argon promotes easy arc starting and stable arc operation. Its lower thermal conductivity promotes the development of axial "spray" transfer in certain forms of GMAW. It is also used in applications where base material distortion must be controlled or where good gap-bridging ability is required. Helium. Unlike argon, helium is lighter than air and has a low density. Like argon, it is chemically inert and does not

react with other elements or compounds. Because of its high thermal conductivity and high ionization potential, more heat is transferred to the base material, thus enhancing the penetration characteristics of the arc. In many applications, it also allows higher weld travel speeds to be obtained. Because of its higher cost, helium is frequently combined with argon or argon mixtures to enhance the overall performance of the blend while minimizing its cost. Oxygen combines with almost all known elements except rare and inert gases; it vigorously support combustion. Small

amounts of oxygen are added to some inert mixtures to improve the stability of the welding arc developed as well as to increase the fluidity of the weld puddle. In the spray-transfer mode of GMAW, small additions of oxygen enhance the range over which this spatterless form of welding can be performed. The droplet size decreases and the number of drops transferred per unit time increases as oxygen is added to the blend. Carbon dioxide is a reactive gas that is commonly used alone in certain types of GMAW. Oxidation of the base material

and any filler electrode occurs readily. Carbon dioxide is added to argon blends to improve are stability, enhance penetration, and improve weld puddle flow characteristics. The higher thermal conductivity of carbon dioxide (because of the dissociation and recombination of its component parts) transfer more heat to the base material than does argon alone. A broader penetration pattern versus argon is obtained; however, base material distortion and lack of gap-bridging ability are possible problems. Hydrogen is the lightest known element and is a flammable gas. Explosive mixtures can be formed when certain

concentration of hydrogen are mixed with oxygen or air. It is added to inert gases to increase the heat input to the base material or for operations involving cutting and gouging. Because some materials are especially sensitive to hydrogenrelated contamination, its use is generally limited to special applications, such as the joining of stainless steels, and to plasma are cutting and gouging. Nitrogen is generally considered to be inert except at high temperatures. At arc welding temperatures, it will react with

some metals (e.g., aluminum, magnesium, steel, and titanium), so it is not used as a primary shielding gas. It can be used with other gases for some welding applications (e.g., copper) and is also widely used in plasma cutting. Shielding Gases for Welding Kevin A. Lyttle, Praxair, Inc.

Shielding Gas Selection In most welding applications, more than one shielding gas or gas blend can be used successfully. For example, there is no one optimal gas blend for joining carbon steels, but a considerable array of mixes are available depending on the specific requirements of the application. For some processes, such as gas-tungsten arc welding (GTAW) and plasma arc welding (PAW), the choices may be somewhat limited by the nature of the electrodes used and the materials being welded. However, for applications involving GMAW, a multitude of blends can be selected from when carbon steels are being joined. Determination of the best blend depends on a number of specific job-related needs. Accuracy of Gas Blends The accuracy with which gases are blended is a function of the way in which they are supplied. If the source of the gas is a high-pressure cylinder, the following generally applies:

1. ±10% RELATIVE, MINOR COMPONENT (REF 3)

2. ±0.5% ABSOLUTE FOR CONCENTRATIONS UP CONCENTRATIONS BETWEEN 5 AND 50% (REF 4)

TO

5%;

±10%

RELATIVE

FOR

For example, the mix accuracy of an Ar-2O2 blend would be Ar/1.8-2.2O2 (method 1) or Ar/1.5-2.5O2 (method 2). A blend of Ar-25CO2 would yield Ar/22.5-27.5CO2 by their method of calculation. When gas cylinders are properly filled with the appropriate blend, the components of that mixture will not separate unless the temperature of the environment is reduced far below normal working temperatures. If the gases are supplied from a liquid source, such as a bulk tank, the accuracy of the blend is a function of the mixing equipment used, but most likely falls within the ±10% minor component range. Shielding Gases for GMAW By far, the largest number of gas blends have been developed for GMAW, especially for joining carbon steel. These can be roughly divided into four categories: pure gases, argon-oxygen mixes, argon/carbon dioxide mixes, and three-part gas blends composed of either argon, helium, oxygen, carbon dioxide, or hydrogen. Table 3 contains suggestions for shielding gas selection based on material type, thickness, and mode of metal transfer.

TABLE 3 RECOMMENDED SHIELDING GAS SELECTION FOR GMAW

MATERIAL CARBON STEEL

THICKNESS TRANSFER MODE MM IN. 0.125

SHORT CIRCUITING

AR-15CO2 AR-25CO2 CO2

GLOBULAR

AR-25CO2 CO2

CONVENTIONAL SPRAY ARC

AR-1O2 AR-2O2 AR-5CO2

ADVANTAGES AND LIMITATIONS GOOD PENETRATION AND DISTORTION CONTROL TO REDUCE POTENTIAL BURNTHROUGH HIGHER DEPOSITION RATES WITHOUT BURNTHROUGH: MINIMUM DISTORTION AND SPATTER; GOOD PUDDLE CONTROL FOR OUT-OF-POSITION WELDING HIGH WELDING SPEEDS, GOOD PENETRATION AND PUDDLE CONTROL; APPLICABLE LOT OUTOF-POSITION WELDS SUITABLE FOR HIGHCURRENT AND HIGHSPEED WELDING; DEEP PENETRATION AND FAST TRAVEL SPEEDS, BUT WITH GREATER BURNTHROUGH POTENTIAL GOOD ARC STABILITY; PRODUCES A MORE FLUID PUDDLE AS O2

PULSED SPRAY

ALLOY STEEL

ALL SIZES

STAINLESS ALL SIZES STEEL, COPPER, NICKEL, AND CU-NI ALLOYS

AR-8CO2 AR-10CO2 AR-15CO2 AR-CO2-O2 BLENDS ARGON-5CO2 AR-HE-CO2 BLENDS AR-CO2-O2 BLENDS

SHORT CIRCUITING

AR-8CO2 AR-15CO2 AR-CO2-O2 BLEND

SPRAY ARC (HIGH-CURRENT DENSITY AND ROTATIONAL)

AR-2O2 AR-5O2 AR-CO2-O2 BLENDS AR-HE-CO2 BLENDS

PULSED SPRAY

AR-5CO2 AR-8CO2 AR-2O2

SHORTCIRCUITING TRANSFER

AR-HE-CO2 BLENDS HE-AR-CO2 BLENDS AR-1O2 AR-2O2

INCREASES; GOOD COALESCENCE AND BEAD CONTOUR, GOOD WELD APPEARANCE AND PUDDLE CONTROL USED FOR BOTH GAGE AND OUT-OF-POSITION WELDMENTS; ACHIEVES GOOD PULSED SPRAY STABILITY OVER A WIDE RANGE OF ARC CHARACTERISTICS AND DEPOSITION RANGES HIGH WELDING SPEEDS; GOOD PENETRATION AND PUDDLE CONTROL; APPLICABLE FOR OUTOF-POSITION WELDS; SUITABLE FOR HIGHCURRENT AND HIGHSPEED WELDING REDUCES UNDERCUTTING; HIGHER DEPOSITION RATES AND IMPROVED BEAD WETTING; DEEP PENETRATION AND GOOD MECHANICAL PROPERTIES USED FOR BOTH LIGHTGAGE AND HEAVY OUT-OF-POSITION WELDMENTS; ACHIEVES GOOD PULSED SPRAY STABILITY OVER A WIDE RANGE OF ARC CHARACTERISTICS AND DEPOSITION RANGES LOW CO2 CONTENTS IN HELIUM MIX MINIMIZE CARBON PICKUP, WHICH CAN CAUSE INTERGRANULAR CORROSION WITH SOME ALLOYS; HELIUM IMPROVES WETTING ACTION; CO2 CONTENTS >5% SHOULD BE USED WITH CAUTION ON

ALUMINUM, TITANIUM, AND OTHER REACTIVE METALS

≤ 13

>13

≤

>

1 2

1 2

SPRAY ARC

AR-HE-CO2 BLENDS AR-1O2 AR-2O2

PULSED SPRAY

AR-HECO2BLENDS AR-1O2 AR-2O2

SPRAY ARC

ARGON

SPRAY ARC

75HE-25AR 50HE-50AR

SPRAY ARC

HELIUM 50HE-25AR

PULSED SPRAY

ARGON

SOME ALLOYS; APPLICABLE FOR ALL POSITION WELDING GOOD ARC STABILITY; PRODUCES A FLUID BUT CONTROLLABLE WELD PUDDLE; GOOD COALESCENCE AND BEAD CONTOUR; MINIMIZES UNDERCUTTING ON HEAVIER THICKNESSES USED FOR BOTH LIGHTGAGE AND HEAVY OUT-OF-POSITION WELDMENTS; ACHIEVES GOOD PULSED SPRAY STABILITY OVER A WIDE RANGE OF ARC CHARACTERISTICS AND DEPOSITION RANGES BEST METAL TRANSFER, ARC STABILITY, AND PLATE CLEANING; LITTLE OR NO SPATTER; REMOVES OXIDES WHEN USED WITH DCEP (REVERSE POLARITY) HIGH HEAT INPUT; PRODUCES FLUID PUDDLE, FLAT BEAD CONTOUR, AND DEEP PENETRATION; MINIMIZES POROSITY HIGH HEAT INPUT; GOOD FOR MECHANIZED WELDING AND OVERHEAD; APPLICABLE TO HEAVY SECTION WELDING GOOD WETTING; GOOD PUDDLE CONTROL

Argon. Pure argon is generally used on nonferrous base metal, such as aluminum, nickel, copper, and magnesium alloys,

and on reactive metals, such as titanium. Argon provides excellent arc stability, penetration, and bead profile when joining these materials. Its low ionization potential results in easy arc starting. Argon produces a constricted arc column with high current density, which concentrates the arc energy over a small area; deep, fingerlike penetration results. Carbon Dioxide. A reactive gas, carbon dioxide is generally used only for joining carbon steel. It is readily available and

relatively inexpensive. Because CO2 will not support spray transfer, deposition efficiency is lower and spatter and fume

levels are higher than with argon blends. Weld bead surfaces are more oxidized and irregular in shape. The higher ionization potential of CO2 and its characteristic dissociation upon heating provide greater weld fusion and penetration while still achieving acceptable mechanical properties. Helium. Because of its higher thermal conductivity, helium can provide additional heat input to the base material while

still maintaining an inert atmosphere. Wetting action, depth of fusion, and travel speed can be improved over comparable argon levels. This advantage is most frequently utilized in the welding of heavier sections of aluminum, magnesium, and copper alloys. Argon-Oxygen. The addition of a small amount of oxygen to a argon greatly stabilizes the welding arc, increases the

filler metal droplet rate, lowers the spray transition current, and influences bead shape. The weld pool is more fluid and stays molten longer, allowing the metal to flow out toward the edges of the weld (Fig. 1a).

FIG. 1 EFFECT OF SHIELDING GAS BLENDS ON WELD PROFILE USING DIRECT CURRENT ELECTRODE POSITIVE (DCEP). (A) ARGON VERSUS ARGON-OXYGEN. (B) CARBON DIOXIDE VERSUS ARGON/CARBON DIOXIDE. (C) HELIUM VERSUS ARGON-HELIUM. SOURCE: REF 5

The most common blends contain 1, 2, 5, or 8% O2 in argon. Increasing oxygen improves arc stability and makes higher travel speeds possible by enhanced puddle fluidity. Some increased alloy loss and a greater chance of undercut occur as the oxygen level is increased, especially beyond 5%. Argon/carbon dioxide blends are primarily used for carbon and low-alloy steels and have limited use for stainless

steels. The addition of CO2 to argon produces results similar to the addition of oxygen, but is also broadens the penetration pattern as the CO2 content is increased (Fig. 1b). Above a range of 18 to 20% CO2, spraylike transfer can no longer be obtained; short-circuiting/globular transfer with somewhat increased spatter levels is found from this point up to approximately 50% CO2 in argon. The most common blends for spray transfer are argon plus 5, 8, 10, or 13 to 18% CO2. With increased CO2 content, the more fluid weld puddle permits higher weld travel speeds. Mixes with higher carbon dioxide levels can also be used for short-circuiting transfer--commonly, argon plus 20 or 25% CO2. Mixtures in this range provide an optimum droplet frequency for minimum spatter when small-diameter (0.9 and 1.2 mm, or 0.035 and 0.045 in.) wire is used. Argon-Helium. Helium is often mixed with argon to obtain the advantages of both gases. These blends are primarily

used for nonferrous base materials, such as aluminum, copper, and nickel alloys. Helium increases the heat input to the base material and thus is used for joining thick, thermally conductive plates. As the helium percentage increases, the arc voltage, spatter, and weld width-to-depth ratio increase (Fig. 1c). The most common blends contain 25, 50, or 75% He in argon. The highest percentage of helium is used for joining thick (>50 mm, or 2 in.) plate material, especially aluminum and copper. Higher travel speeds can be obtained using heliumenhanced blends. Argon/Oxygen/Carbon Dioxide. Mixtures containing three gas components are versatile because of their ability to

operate in short-circuiting, globular, and spray-transfer modes. These blends are generally proprietary, and manufacturers' recommendations should be followed for their proper use. Blends of argon, carbon dioxide, and oxygen are generally used to join carbon and alloy steels.

Argon/Helium/Carbon Dioxide. Helium and carbon dioxide additions to argon increase the heat input to the base

metal, which improves wetting, puddle fluidity, and weld bead profile. With helium plus CO2 additions less than 40%, good spray transfer is obtained for carbon and low-alloy steel welding. Some increase in tolerance of base material surface contamination is also noted. When the helium content exceeds 50 to 60%, transfer is restricted to short-circuiting and globular. Blends in which the CO2 content is relatively low ( ≤ 5%) are generally used for the joining of stainless steels without any loss in corrosion resistance. Argon/Carbon Dioxide/Hydrogen. Three-part blends of this design are intended for the joining of austenitic stainless

in the spray or short-circuiting transfer modes. Because of the addition of hydrogen, these blends should not be used for carbon steel. The carbon dioxide and hydrogen increase the heat input to the base material and improve bead shape characteristics, as well as promote higher welding travel speeds. Shielding Gas Selection for FCAW Carbon Dioxide. The majority of large-diameter (>1.6 mm, or

1 in.) wires that use a shielding gas use carbon dioxide. 16

Some smaller diameter wires are formulated to operate in 100% CO2. The arcs are generally stable and provide a globular transfer over the usable operating range. Good performance over rust and mill scale on the plate surface is obtained with these large-diameter wires and CO2 shielding. Argon/Carbon Dioxide. A significant number of small-diameter ( ≤ 1.6 mm, or

1 in.) cored wires are shielded with the 16

blends of argon with 15 to 50% CO2. These blends provide better out-of-position weld puddle control versus that of CO2. To obtain the best performance from a particular cored wire, check the manufacturer's product literature for the recommended gas blend. Shielding Gas Selection for GTAW Argon. The most commonly used gas for GTAW, argon exhibits low thermal conductivity, which produces a narrow,

constricted arc column; this allows greater variations in arc length with minimal influence on arc power or weld bead shape. Its low ionization potential provides good arc starting characteristics and good arc stability using the direct current electrode negative (DCEN) power connection plus superior arc cleaning action and bead appearance when ac power is used. Argon is the most commonly selected gas for DCEN welding of most materials and ac manual welding of aluminum. Helium. The high thermal conductivity and ionization potential of helium make it suitable for the high-current joining of

heavy sections of heat-conductive materials such as aluminum. Helium increases the penetration of the weld as well as its width. It also allows the use of higher weld travel speeds. Argon-Helium. Blends of argon and helium are selected to increase the heat input to the base material while maintaining

favorable arc stability and superior arc starting characteristics. Blends of 25, 50, and 75% He in argon are commonly used. Argon-Hydrogen. Hydrogen is added to argon to enhance its thermal properties. The slightly reducing atmosphere

improves weld puddle wetting and reduces some surface oxides to produce a cleaner weld surface. To minimize problems associated with arc starting, additions of hydrogen are generally limited to 5 to 15%. These blends are primarily used to join some stainless steels, nickel, and nickel alloys. These mixtures should not be used to join alloy steels; delayed weld cracking may result. Argon/2-5H2 is used in manual welding applications on materials thicker than 1.6 mm (

1 in.). Additions of 10 to 15% 16

H2 are used in mechanized applications, such as those found in the manufacture of stainless steel tubing. Caution: Special safety precautions are required when mixing argon and hydrogen. Do not attempt to mix argon and hydrogen from separate cylinders. Always purchase ready-mixed hydrogen blends from a qualified supplier.

Shielding Gas Selection for PAW The physical configuration of the PAW system requires the use of two gases: a "plasma" or orifice gas and a shielding gas. The primary role of the plasma gas, which exits the torch through the center orifice, is to control arc characteristics and shield the electrode. The shielding gas, introduced around the periphery of the arc, shields or protects the weld area. In many applications, the shielding gas is also partially ionized to enhance the performance of the plasma gas. Low-Current (1.6 mm, or

1 in.) wires that use a shielding gas use carbon dioxide. 16

Some smaller diameter wires are formulated to operate in 100% CO2. The arcs are generally stable and provide a globular transfer over the usable operating range. Good performance over rust and mill scale on the plate surface is obtained with these large-diameter wires and CO2 shielding. Argon/Carbon Dioxide. A significant number of small-diameter ( ≤ 1.6 mm, or

1 in.) cored wires are shielded with the 16

blends of argon with 15 to 50% CO2. These blends provide better out-of-position weld puddle control versus that of CO2. To obtain the best performance from a particular cored wire, check the manufacturer's product literature for the recommended gas blend. Shielding Gas Selection for GTAW Argon. The most commonly used gas for GTAW, argon exhibits low thermal conductivity, which produces a narrow,

constricted arc column; this allows greater variations in arc length with minimal influence on arc power or weld bead shape. Its low ionization potential provides good arc starting characteristics and good arc stability using the direct current electrode negative (DCEN) power connection plus superior arc cleaning action and bead appearance when ac power is used. Argon is the most commonly selected gas for DCEN welding of most materials and ac manual welding of aluminum. Helium. The high thermal conductivity and ionization potential of helium make it suitable for the high-current joining of

heavy sections of heat-conductive materials such as aluminum. Helium increases the penetration of the weld as well as its width. It also allows the use of higher weld travel speeds. Argon-Helium. Blends of argon and helium are selected to increase the heat input to the base material while maintaining

favorable arc stability and superior arc starting characteristics. Blends of 25, 50, and 75% He in argon are commonly used. Argon-Hydrogen. Hydrogen is added to argon to enhance its thermal properties. The slightly reducing atmosphere

improves weld puddle wetting and reduces some surface oxides to produce a cleaner weld surface. To minimize problems associated with arc starting, additions of hydrogen are generally limited to 5 to 15%. These blends are primarily used to join some stainless steels, nickel, and nickel alloys. These mixtures should not be used to join alloy steels; delayed weld cracking may result. Argon/2-5H2 is used in manual welding applications on materials thicker than 1.6 mm (

1 in.). Additions of 10 to 15% 16

H2 are used in mechanized applications, such as those found in the manufacture of stainless steel tubing. Caution: Special safety precautions are required when mixing argon and hydrogen. Do not attempt to mix argon and hydrogen from separate cylinders. Always purchase ready-mixed hydrogen blends from a qualified supplier.

Shielding Gas Selection for PAW The physical configuration of the PAW system requires the use of two gases: a "plasma" or orifice gas and a shielding gas. The primary role of the plasma gas, which exits the torch through the center orifice, is to control arc characteristics and shield the electrode. The shielding gas, introduced around the periphery of the arc, shields or protects the weld area. In many applications, the shielding gas is also partially ionized to enhance the performance of the plasma gas. Low-Current ( γSL > γLV

(EQ 2)

For soldering, the vapor phase will be replaced in nearly all processes by flux (that is, γsv is replaced by γsf and γlv is replaced by γlf). For optimum wetting, the contact angle, θ, must be minimized and, therefore, γsv must be maximized. Oxides and contamination lower the value of γsv (or γsf). A function of the flux is to remove or disperse oxides and contaminants, thereby increasingly γsv. An important function of flux is to lower the value of γlf (or γlv). The flux also maintains a local environment around the surfaces being joined, protecting them from reoxidation during soldering, as well as enhancing heat transfer from the heating source to the substrate and the solder. Equation 1 shows that wetting is also enhanced by minimizing γlv, a value commonly referred to as the surface tension of the solder. Although γlv is a function of the solder composition, the flux covering the liquid solder, and the temperature, it is normally close to 0.4 J/m2 (4 × 10-5 Btu/ft2). A typical example of nonwetting occurs when a liquid metal droplet is placed on a nonmetallic surface, such as copper oxide. Because there are no metallic bonds at the oxide surface, the liquid metal will have little tendency to interact with that surface. The metal droplet will behave as though it is repelled, and will try to ball up to minimize the area of contact with the nonmetallic surface. One method for quantifying wettability is to observe the behavior of solder droplets on the metal surface. The contact angle formed between the surface of the liquid solder and the surface of the solid can be

measured (Fig. 5). If the solder droplet forms an angle, θ, of less than 90°, then the solder is said to wet the metal. If the angle is greater than 90°, the solder is nonwetting. This is one method used to qualitatively assess solderability. The formation of intermetallic compounds can affect Young's equation because it alters the value of γsl. The influence of this term has been recognized (Ref 6), but the magnitude has not yet been calculated quantitatively. Surfaces of a metal crystal have a higher free energy than the bulk of the crystal because of the existence of unsatisfied metallic bonds. When a liquid metal or alloy is brought into contact with this solid surface, the liquid will proceed to interact with the solid to satisfy those dangling surface bonds and thus reduce the surface free energy of the system. To the extent that a liquid solder can satisfy these surface bonds, it will wet the solid metal and spread across its surface. If no exchange of bonding energy takes place, then wetting will not occur and the solder will tend to ball up on the surface of the solid, minimizing the area of contact between liquid and solid. Information about selecting solders for specific applications is provided in the article "General Soldering" in this Volume.


3. R.J. KLEIN WASSINK, SOLDERING IN ELECTRONICS, ELECTROCHEMICAL PUBLICATIONS LTD., 1989 6. A.D. ROMIG, JR., Y.A. CHANG, J.J. STEPHENS, ET AL., PHYSICAL METALLURGY OF SOLDERSUBSTRATE REACTIONS, SOLDER MECHANICS, THE MINERALS, METALS, & MATERIALS SOCIETY, 1991, P 30-32 Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

Guidelines for Flux Selection A flux promotes solder wetting of the base materials by: • • •

REMOVING TARNISH FILMS FROM PRECLEANED SURFACES PREVENTING OXIDATION DURING THE SOLDERING OPERATION LOWERING THE SURFACE TENSION OF THE SOLDER

Successful establishment of a solder joint requires that the liquid solder make contact with the metal to which it is to be joined, so that wetting can be initiated. Unfortunately, almost all of the metals involved in soldering are oxidized during exposure to elevated temperatures in air. This prevents metal-to-metal contact, as well as the wetting and formation of a metallurgical bond, unless the oxides are removed. Fluxes are used in soldering to remove such films and to protect the surfaces against reoxidation during soldering. In addition to cleaning metal surfaces, a flux also: • • •

AIDS IN THE WETTING PROCESS IMPROVES HEAT TRANSFER (ESPECIALLY IN SOLDER PASTES) CARRIES OXIDES AND SOLID DEBRIS AWAY FROM THE JOINT

Fluxes are chemical (liquid, solid, or gaseous materials) that remove oxide layers from the base metal and solder. When heated, fluxes either promote or accelerate the wetting of metals by solder. Flux selection usually depends on the ease with which a material can be soldered. Rosin fluxes are used with solderable base metals or with metals that are precoated

with a solderable base metal or with metals that are precoated with a solderable finish. Inorganic fluxes are often used on metals such as stainless steel. Table 1 indicates the relative ease with which a number of alloys and metals can be soldered, based on flux requirements.

TABLE 1 RELATIVE SOLDERABILITY OF SELECTED METALS AND ALLOYS AS A FUNCTION OF FLUX TYPE USED Base Metal, Alloy, Or Applied Finish

Flux Type ROSIN ORGANIC

INORGANIC

SPECIAL FLUX AND SOLDER

Aluminum Aluminum-Bronze Beryllium Beryllium-Copper Brass Cadmium Cast Iron Chromium Copper Copper-Chromium Copper-Nickel Copper-Silicon Gold Inconel Lead Magnesium Manganese-Bronze (High Tensile) Monel Nickel Nickel-Iron Nichrome Palladium Platinum Rhodium Silver Stainless Steel Steel Tin Tin-Bronze Tin-Lead Tin-Nickel Tin-Zinc Titanium Zinc Zinc Die Castings

... ... ... X X X ... ... X ... X ... X ... X ... ... ... ... ... ... X X ... X ... ... X X X ... X ... ... ...

... ... ... X X X ... ... X X X X X ... X ... ... X X X ... X X X X X X X X X X X ... X ...

X X ... ... ... ... X ... ... ... ... ... ... X ... ... ... ... ... ... X ... ... ... ... ... ... ... ... ... ... ... ... ... ...

... ... ... X X X ... ... X ... X ... X ... X ... ... X X X ... X X ... X X ... X X X X X ... X ...

Soldering Not Recommended(A)

... ... X ... ... ... ... X ... ... ... ... ... ... ... X X ... ... ... ... ... ... ... ... ... ... ... ... ... ... ... X ... X

(A) WITH PROPER PROCEDURES, SUCH AS PRECOATING, MOST METALS CAN BE SOLDERED Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

Types of Fluxes Fluxes can be categorized by their chemical makeup (Ref 2, 3, 9, 10, 11): • • • •

ROSIN-BASE FLUXES ORGANIC FLUXES INORGANIC FLUXES SYNTHETICALLY ACTIVATED FLUXES (RESINS)

These materials vary in their activity (that is, aggressiveness). Rosin-base fluxes are very mild, whereas inorganic fluxes can be extremely active and corrosive. Because of the potential for corrosion by inorganic fluxes in electronic assemblies, the rosin-base fluxes (with their low flux activity) are favored by the electronics industry. Rosin Fluxes. Water-white rosin dissolved in a suitable organic solvent is the closest model of a flux with noncorrosive

residue. Rosin fluxes possess important physical and chemical properties that make them particularly suitable for use in the electrical industry. They are solid and inactive chemically at room temperature, but at soldering temperatures they gain sufficient activity to remove weakly adherent oxides from the noble metals gold and silver, as well as copper. Nonactivated Rosin Flux. Although its noncorrosive nature has led to its widespread use in microelectronics, rosin-

base flux is lacking in chemical activity, which essentially limits its use to precleaned parts and to only a few metals that do not have adherent oxides. The active constituent, abietic acid (C20H30O2), with a melting point of 173 °C (343 °F), becomes mildly active at soldering temperatures ranging from 177 to 316 °C (350 to 601 °F). The residue is hard, nonhygroscopic, electrically nonconductive, and noncorrosive. Mildly Activated Rosin Flux. The very low chemical activity of rosin-base flux can be countered by the addition of

activators, usually organic acids or amines. The end product, rosin mildly activated (RMA) flux, is still essentially noncorrosive, but is sufficiently active to remove oxides more reliably than the rosin-base flux. In addition, the residues are noncorrosive. However, care must be exercised in cases of high-reliability applications. It is advisable to specify that these fluxes be halide-free, in order to ensure that specific safety requirements are met in critical components. Even so, cleaning is now required to remove the flux residue. Both polar and nonpolar solvents have been used. These RMA fluxes are preferred for military, telephone, and other high-reliability electronic products. Activated rosin fluxes were developed to provide more chemically active fluxes for mass-produced electronics, such as

packaged components. Most mass-produced electronics are manufactured using RMA fluxes. The use of chlorides in these fluxes requires effective cleaning after soldering to prevent corrosion and electrical leakage, because the presence of chloride ions in flux residue makes it conductive. Because of the chemistry of the activators, a double solvent cleaning is normally required. Activated rosin fluxes are widely used in commercial electronics and in high-reliability applications where the residue can be completely removed after soldering. Inorganic Fluxes. Inorganic acids and salts that are highly corrosive and extremely active compose this class of fluxes.

In many instances, they are not acceptable for use on electronics, but are suitable for plumbing and industrial applications. Inorganic fluxes do enable the soldering of ferrous alloys and high-nickel alloys used in electronic packages and hermetic enclosures. The difficulty of chloride ion removal has led to the gradual abandonment of inorganic fluxes, even for tinning purposes. Inorganic fluxes are used to optimum advantage where conditions require rapid and highly active fluxing action. They can be applied as solutions, pastes, or dry salts in many general soldering applications. They function equally well with torch, oven, resistance, or induction soldering methods, because they neither char nor burn. These fluxes can be formulated to provide stability over the entire soldering temperature range. One distinct disadvantage of inorganic fluxes is that their residue remains chemically active after soldering. If this residue is not removed, then severe corrosion can occur at the joint. Adjoining areas can also be attacked by residues from the spray of flux and from flux vapors. Fluxes that contain ammonium salts can cause stress-corrosion cracking (SCC) when soldering brass, as well as most iron-nickel alloys.

Organic acid fluxes, although less active than inorganic fluxes, are effective at soldering temperatures ranging from 90

to 320 °C (195 to 610 °F). They consist of organic acids and bases and, often, hydrohalides. They are active at soldering temperatures, but the period of activity is short because of their susceptibility to thermal decomposition. Their tendency to volatilize, char, or burn when heated limits their use with torch or flame heating. When these fluxes are properly used, their residues are relatively inert and can be removed with water. Organic acid fluxes are particularly useful in applications where controlled quantities of flux can be applied and where sufficient heat can be used to fully decompose or volatilize the corrosive constituents. Caution is necessary to prevent undecomposed flux from spreading to insulating sleeving. Care must also be taken when soldering in closed systems where corrosive fumes can condense on critical parts of the assembly.


2. H.H. MANKO, SOLDERS AND SOLDERING, 2ND ED., MCGRAW-HILL, 1964 3. R.J. KLEIN WASSINK, SOLDERING IN ELECTRONICS, ELECTROCHEMICAL PUBLICATIONS LTD., 1989 9. J.F. SHIPLEY, INFLUENCE OF FLUX, SUBSTRATE AND SOLDER COMPOSITION ON SOLDER WETTING, WELD. J., VOL 54 (NO. 10), OCT 1975, P 357S-362S 10. H.H. MANKO, SOLDERING FLUXES--PAST AND PRESENT, WELD. J., VOL 52 (NO. 3), MARCH 1973, P 163-166 11. R.W. WOODGATE, THE HANDBOOK OF MACHINE SOLDERING, 2ND ED., JOHN WILEY & SONS, 1988 Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

Flux Evaluation Several tests are used to evaluate the relative activity of fluxes. The most common are the copper mirror test, the halide test, a surface conductivity test, and a test based on the resistivity of a water extract present in the flux. In all cases, a measure of the chemical activity of the flux is obtained. Because the fluxes all depend on other properties (for example, viscosity, ease of removal, and thermal behavior), chemical activity is only one factor in flux selection for a given application. Information about selecting fluxes for specific applications is provided in the article "General Soldering" in this Volume. Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

Joint Design Joints should be designed to fulfill the requirements of the finished assembly, as well as to permit the application of the flux and solder. Joint design should maintain proper clearance during heating and upon solidification of the filler metal. Special fixtures may be necessary or the units can be crimped, clinched, wrapped, or otherwise held together. The selection of a joint design for a specific application will primarily depend on the service requirements of the assembly. It may also depend on such factors as the heating method to be used, the fabrication techniques utilized prior to soldering, the number of items to be soldered, and the method used to apply the solder. In general, solders have low

strength when compared with the metals for which they are used to join. Therefore, the soldered joint should be designed to avoid dependence on solder strength. The necessary strength can be provided by shaping the parts to be joined so that they engage or interlock, requiring the solder only to bond, seal, and stiffen the assembly. Figure 6 shows joint designs commonly used for soldering applications. The lap joint and the lock seam joint are used when soldering sheets. Lap joints are also applied to join pipes. Actually, the lap type of joint should be used whenever possible because it offers the best chance to obtain joints with maximum strength. It should also be used whenever a seal is required. Butt joints should be avoided whenever possible.

FIG. 6 BASIC JOINT CONFIGURATIONS USED FOR SOLDERING APPLICATIONS. (A) LAP JOINT. (B) LOCK SEAM JOINT. (C) BUTT JOINT. (D) PIPE JOINT

An important factor in joint design is the manner in which the solder will be applied to the joint. The designer must consider the number of joints per assembly and the number of assemblies to be manufactured. For limited production using a manual soldering process, the solder can be face-fed into the joint with few problems. However, for a large production lot of assemblies containing multiple joints, an automated process such as condensation/vapor-phase reflow soldering can be advantageous (Ref 12, 13, 14). In this case, the design must provide for accessible joints that are suitable for automated fluxing, soldering, and cleaning.

There should be enough clearance between the parts being joined to allow the solder to be drawn into the space by capillary action, but not so much that the solder is unable to fill the gap. Joint clearances ranging from 0.075 to 0.150 mm (0.0030 to 0.059 in.) are preferred for optimum strength, but variations are allowed in specific cases. It is often necessary to fabricate sample parts and to test the joints to ensure their producibility and capability regarding strength properties (Ref 2, 3, 10, 15, 16, 17). Other commonly used soldered joints are shown in Fig. 7; self-jigging joints are shown in Fig. 8. The most definitive work available on the loading of soldered joints is that used for the load-carrying capabilities of copper tube with sleevetype joints or fittings. Conservative joint design requires that only 50% of the joint be considered filled. Under normal conditions, it would be considered good soldering practice if 70% or more of the joint consists of sound joint material (Ref 2, 3, 4, 5, 6, 11, 15, 18). Additional information about joint configurations is provided in the article "General Soldering" in this Volume.

FIG. 7 JOINT DESIGNS FREQUENTLY USED IN SOLDERING OPERATIONS. SOURCE: REF 5

FIG. 8 METHODS THAT CAN BE USED TO MAKE SOLDER JOINTS SELF-JIGGING. SOURCE: REF 5


2. H.H. MANKO, SOLDERS AND SOLDERING, 2ND ED., MCGRAW-HILL, 1964 3. R.J. KLEIN WASSINK, SOLDERING IN ELECTRONICS, ELECTROCHEMICAL PUBLICATIONS LTD., 1989 4. L.P. LAMBERT, SOLDERING FOR ELECTRONIC ASSEMBLIES, MARCEL DEKKER, 1987 5. SOLDERING MANUAL, AWS, 1978 6. A.D. ROMIG, JR., Y.A. CHANG, J.J. STEPHENS, ET AL., PHYSICAL METALLURGY OF SOLDERSUBSTRATE REACTIONS, SOLDER MECHANICS, THE MINERALS, METALS, & MATERIALS SOCIETY, 1991, P 30-32 10. H.H. MANKO, SOLDERING FLUXES--PAST AND PRESENT, WELD. J., VOL 52 (NO. 3), MARCH 1973, P 163-166 11. R.W. WOODGATE, THE HANDBOOK OF MACHINE SOLDERING, 2ND ED., JOHN WILEY & SONS, 1988 12. T. THOMPSON, CONDENSATION/VAPOR-PHASE REFLOW SOLDERING, ASS. ENG., JUNE 1977, P 44-47 13. T.Y. CHU, A GENERAL REVIEW OF MASS SOLDERING METHODS, INS. CIRC., NOV 1976 14. R.C. PFAHL, JR., J.C. MOLLENDORF, AND T.Y. CHU, CONDENSATION SOLDERING, WELD. J., VOL 54 (NO. 1), 1975 15. E. LIEBERMAN, MODERN SOLDERING AND BRAZING TECHNIQUES, BUSINESS NEWS PUBLISHING CO., 1988 16. "SOLDER METAL SPECIFICATION," ASTM B 32, ANNUAL BOOK OF ASTM STANDARDS, PART B, ASTM 17. D.R. FREAR, W.B. JONES, AND K.R. KINSMAN, ED., SOLDER MECHANICS, TMS, 1991, P 30-41 18. PROC. IEEE 41ST ELECT. COMP. & TECH. CONF., IEEE, 1991 Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

Precleaning and Surface Preparation An unclean surface will not permit the solder to flow, which makes soldering either difficult or impossible and contributes to the formation of a poor joint. Materials such as oil, grease, paint, pencil markings, drawing, and cutting lubricants, general atmospheric dirt, oxide, and rust films must be removed before soldering. The importance of cleanliness cannot be overemphasized in ensuring sound soldered joints. Fluxing alone cannot substitute for adequate precleaning (Ref 19). Because cleaning methods are often designed for a specific soldering operation, their suitability for a particular application should be investigated thoroughly. Degreasing. Either solvent or alkaline degreasing is recommended for the cleaning of oily or greasy surfaces. Of the

solvent degreasing methods, the vapor condensation of the trichloroethylene-type solvents probably minimizes residual film on the surface. In the absence of vapor degreasing apparatus, immersion in liquid solvents or in detergent solutions is a suitable procedure. Hot alkali detergents are widely used for degreasing. All cleaning solutions must be thoroughly removed before soldering. Residues from hard-water rinses can interfere with soldering. Material compatibility is a prime consideration. Pickling. The purpose of pickling, or acid cleaning, is to remove rust, scale and oxides, or sulfides from the metal in order to provide a clean surface for soldering. All of the inorganic acids (that is, hydrochloric, sulfuric, phosphoric, nitric, and hydrofluoric acids), singly or mixed, will fulfill this function, although hydrochloric and sulfuric acids are the most widely used. The pieces should be washed thoroughly in hot water after pickling and then dried as quickly as possible. Mechanical cleaning includes the following methods:

• • • • •

GRIT AND SHOTBLASTING MECHANICAL SANDING AND GRINDING FILING AND HAND SANDING CLEANING WITH STEEL WOOL WIRE BRUSHING AND SCRAPING

For best results, cleaning should extend beyond the joint area. Shot or steel gritblasting is often effective and is preferable to sandblasting because it avoids the embedding of silica particles on the surface, which interferes with the flow of solder. Mechanical cleaning is not recommended for softer metals, such as copper. Precoating. The coating of the base metal surfaces with a more-solderable metal or alloy prior to the soldering operation can facilitate soldering. Coatings of tin, copper, silver, cadmium, iron, nickel, and the alloys of tin-lead, tin-zinc, tincopper, and tin-nickel are used for this purpose. The advantages of precoating include the following: • •

SOLDERING IS MORE RAPID AND UNIFORM. STRONG ACIDIC FLUXES CAN BE AVOIDED DURING SOLDERING.

The precoating of metals that have tenacious oxide films (for example, aluminum, aluminum bronzes, highly alloyed steels, and cast iron) is almost mandatory. The precoating of steel, brass, and copper, although not entirely essential, is of great value in some applications. A number of different methods are used to precoat the metal surfaces. Solder or tin can be applied using a soldering iron or an abrasive wheel, ultrasonic soldering, immersion in molten metal, electrodeposition, or chemical displacement. Hot dipping can be accomplished by fluxing and dipping the parts in molten tin or solder. Often, small parts are initially placed in wire baskets and then cleaned, fluxed, and dipped in the molten metal. Finally, they are centrifuged to remove the excess metal. Coating by hot dipping is applicable to carbon steel, alloy steel, cast iron, copper, and certain copper alloys. Prolonged immersion in molten tin or solder should be avoided to prevent the excessive formation of intermetallic compounds at the coating/base-metal interface. Electrodeposition. Precoating by electrodeposition can be accomplished in stationary tanks, in conveyorized plating units, or in barrels. This method is applicable to all steels, copper alloys, and nickel alloys. The coating metals are not limited to tin and solder. Copper, cadmium, silver, precious metals, nickel, iron, and alloy platings such as tin-copper, tinzinc, and tin-nickel are also commonly used.

Certain combinations of electrodeposited metals (that is, duplex coatings), where one metal is plated over another, are becoming more popular as an aid to soldering. A coating of 0.005 mm (0.0002 in.) of copper plus 0.0075 mm (0.0003 in.) of tin is particularly useful for brass. The solderability of aluminum is enhanced by a coating of 0.013 mm (0.0005 in.) of nickel, followed by 0.008 mm (0.0003 in.) of tin or by a combination of zincate (zinc), 0.005 mm (0.0002 in.) copper, and tin. An iron plating followed by tin plating is extremely useful over a cast iron surface. Immersion coatings (that is, chemical displacement coatings) of tin, silver, or nickel can be applied to some of the

common base metals. These coatings are usually very thin and generally have a poor shelf life. The shelf life of a coating is defined as its ability to withstand storage conditions without impairing solderability. Hot tinned and brightened electrotin coatings have an excellent shelf life. Inadequate thicknesses of electrotinned or immersion-tinned coatings have a limited shelf life. Tin or solder coating thicknesses ranging from 0.003 to 0.008 (0.0001 to 0.0003 in.) are recommended to ensure maximum solderability after prolonged storage.


19. J. BROUS, EVALUATION OF POST-SOLDER FLUX REMOVAL, WELD. J., VOL 54 (NO. 12), 1975, P 444S-448S

Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

Soldering in Electronic Applications This section addresses some general considerations in the soldering of electronic devices. More detailed information is provided in the article "Soldering in Electronic Applications," in this Volume. Substrates for Electronic Component Applications A substrate is the platform on which the electrical/electronic circuitry is built. Many different types of substrates are used in the electronics field. Early substrates were sheets of material to which insulated, stand-off terminals were attached by bolting or riveting. The circuitry was then obtained by point-to-point wiring between these terminals. Although this approach is still used, the most common substrate in use today has circuitry that is laid on it in planar form, with the components already attached. Attachment is normally accomplished through the soldering process. The most common form of this substrate is the printed wiring board (PWB). A substrate can actually have numerous forms: • • •

CIRCUITRY CAN BE LAID ON ONE SIDE OF THE SUBSTRATE OR ON BOTH SIDES. CIRCUITRY CAN BE NOT ONLY ON THE OUTER SURFACES OF THE SUBSTRATE, BUT IN LAYERS WITHIN THE SUBSTRATE MATERIAL (MULTILAYERING). THE SUBSTRATE CAN BE RIGID OR FLEXIBLE.

There are many methods of applying the conductive patterns that make up the circuitry and each has its own specific soldering solutions. Many different materials are used for substrates. Substrates used for electrical/electronic applications must be insulators to allow the conductive circuitry to function in the manner of a network made from separate insulated wires. Common substrate materials for this application include ceramics and organic laminates. Metals can also be used as substrates if they are coated with a substantial insulating material that can withstand the rigors of assembly. Ceramic Substrates. Many types of ceramics can be used as substrates, but the most common is a high-fired alumina

that is either 96 or 99% Al2O3. The 96% Al2O3 material is usually employed when thick-film materials are placed on the surface to form the circuitry. The 99% Al2O3 material is most often used for the evaporation or sputtering of thin materials. The materials that are laid down for terminal and component bonding pads are usually gold (or alloys high in gold), silver, and copper. Because these metals are laid down in thin layers, a soldering problem called leaching occurs. Leaching is the migration of the components of a substrate into a working solution. The process must be controlled because soldering results in bonding between the solder and the metals being attached. If the soldering process is not controlled, it can literally dissolve the metallization from the surface of the substrate. A gold-platinum alloy is more resistant to leaching than pure gold. o

Thin-film (10 to 500 nm, or 100 to 5000 A ) deposits are more susceptible to leaching than are thick-film (10 to 25 μm, or 400 to 1000 μin.) deposits, because the thin layer is more susceptible to being consumed by the leaching. Therefore, the time at temperature and the maximum temperature become very important. In leaching, temperature is more important than time. Significantly, more leaching will be generated by a small increase in temperature than would be produced with even a moderate increase in time. Composite Laminates. The most popular composites are the copper-clad laminates used for PWB substrates. The metal cladding is usually copper, but could be nickel or some other metal. Copper cladding can be either a plated material or wrought rolled foil. From a soldering standpoint, the rolled material is preferable, but it is more costly and more difficult to bond to the laminate material. Many PWBs are passed through numerous heat cycles during the manufacturing process. In addition, the use of multilayer designs is increasing dramatically. Plated foil, contaminated with codeposited organics,

can become brittle at elevated temperatures. Hence, foil coppers are available in several classes, depending on their application. Printed Wiring Laminate Materials. The laminate upon which the metal cladding is applied is available in many composite forms. Originally, layers of paper or cloth impregnated with a phenolic resin and designated XP, XXP, or XXXP were used. The latter became the standard in the mid-1950s.

The selection of organic-based laminate materials as electronic substrates has certain limitations, the most common of which is temperature. In addition, soldering must not degrade the material by burning or charring. Laminate selection should depend on the application of the electronic hardware. A glass-epoxy laminate material known as FR-4 is commonly used in consumer and industrial electronics. The "FR" designation indicates that the material is fire resistant. Metal Substrates. There are some applications in which metals can be used as base substrates. Because the presence of

an insulating material is required, the metal must be coated with a suitable material (for example, porcelain enamel). Solder Application Because of the small distances between components, when soldering electronics, it is often necessary to measure carefully the amount of solder that must be applied to form a satisfactory joint. It is not only a matter of providing a sufficient amount of solder for a good fillet, but it is also important to not have excess solder in the joint. This can sometimes be accomplished by plating the area to be soldered to a predetermined solder thickness, so that the solder forms a good fillet when it is reflowed. This is often the method used to accomplish "flip-chip" bonding to produce tape automated bonding (TAB) assemblies (Ref 20). Solder can also be apportioned in proper amounts to individual joints by using preforms. In the case of surface-mount technology (SMT), the solder is often screened onto the substrate in the form of solder paste or it is dispersed in small quantities in precise locations by using tiny nozzles or tubes. Preforms. When a square or round lid is attached to a device, but solder is not desired in the enclosed area, the solder

application is often accomplished by using squares or annular rings that are punched from solder sheet. The solder quantity is usually sufficient to fill the asperities in the surfaces. The thickness and the lateral dimensions are determined by the location and the amount of solder needed to make a satisfactory joint. The solder preform is placed in the space between the parts to be joined, and then this assembly is subjected to sufficient heat to fuse the solder and hermetically seal the assembly. The entire body is then cooled. If the parts were clean and the proper amount of solder was supplied, a good solder joint results. Heat for soldering can be applied in a number of different ways. In automated processing, the parts to be soldered are often passed through the hot zone of a furnace on a belt. This makes it possible to provide adequate preheating of the parts, the proper maximum temperature, and a gentle cooling cycle to minimize stresses. This method is fast and allows easy control of the measured amounts of solder to precise locations in a variety of geometric shapes. However, the cost of the preforms is a drawback. Solder Pastes. The rapid development of SMT processing, which can involve very dense populations of devices and

leads, necessitated a method to precisely place controlled amounts of solder on a substrate, either for lead attachment or for bonding of the device to attain more-effective heat transfer. Solder pastes that are applied by screening or printing through a mask or by application through a fine tube can fulfill these needs. The screening method, although ancient, has been modernized by the development of new high-precision screens, solder pastes, and equipment for screening and curing the screened materials. Solder paste also can be applied through a metal mask, or stencil. Although the location of the solder is not as precise as that obtained with screening, the metal masks are more durable and provide better control of solder thickness as it is applied. Another paste placement method is to use fine tubes or nozzles to deliver the paste to precise locations. This method is ideal for placing discrete dots of paste, rather than covering large areas.

Paste Composition. A solder paste must contain solder particles of the desired composition to yield the melting and

joining requirements of the application. The metal particles must be sized to provide good screenability and to avoid excessive oxide content. The shape of the particles is important. Spherical particles are lower in oxide content because there is less surface area and they are able to flow more freely in screening applications. A paste must also contain a vehicle to control its rheological, or flow, properties; its drying behavior; its surface activity (wetting behavior); and the solubility of the residues left after reflow. The vehicle usually consists of a rosin (pine derivative) or a synthetic resin, solvents to adjust drying time and temperature, modifiers to control rheological behavior, and an activator, or flux, to remove oxides from the solder and to protect the liquefied solder from oxidation. Paste Properties. The paste must set and hold its shape and location after screening so that subsequent processing

maintains the solder in the desired locations and amounts. The assembly is then subjected to heat and the solder is melted and joined to the part being soldered. This reflow can be conducted either as the final step in soldering or as a separate step. The parts must generally be cleaned after soldering. Solder resists enable control over the PWB areas that will be wet by the solder. The resist, or mask, is an organic

coating placed selectively on the board. Only the uncoated areas are exposed to and are wettable by the solder. The higher interconnect densities now being attained by SMT processing make solder resists a valuable method. Solder resists can be either permanent or temporary, depending on the application. The permanent resists must resist not only the solder, but any fluxes or solvents applied later in the manufacturing sequence. Compatibility with the process and the circuit materials is a consideration, as is the ability to withstand soldering temperatures. In addition, the solder resist offers the circuits some protection against dirt and contamination. Commonly used resist materials are polystyrene, epoxides, and acrylics, although more-costly materials, such as polyimides, are sometimes used because of their greater thermal resistance. Screened Materials. Solder masks are commonly applied by screening. The films are subsequently dried and cured

thermally. The elevated temperatures involved in the curing cycle can cause oxidation of the conductors on the PWB, resulting in a loss of solderability. Lower curing temperatures are desirable. Some of the resists are cured by exposure to ultraviolet light, rather than by heating. This approach is very fast, involves less emission of volatile pollutants, and does not cause oxidation of the metal surfaces. Photographic Films. A second method of applying solder resists is to use sheets of photosensitive material laminated to

the surface of the circuit. Although these are more expensive than screened materials, the photographic films enable much higher pattern resolution. Dry film resists are much stronger and more resistant to damage during soldering and cleaning operations. The use of these photosensitive materials increases the complexity of the process. Temporary solder masks are often used as production aids, offering local protection from solder at locations such as: • • •

CONTACT PADS THAT HAVE BEEN GOLD PLATED PLATED-THROUGH HOLES THAT MUST BE SOLDERED IN A LATER OPERATION PADS THAT MUST BE KEPT FREE OF SOLDER

Adhesive solder-mask tapes that specifically protect edge connectors on PWBs are also available, and numerous mechanical approaches to limiting solder contact exists (Ref 18).


18. PROC. IEEE 41ST ELECT. COMP. & TECH. CONF., IEEE, 1991 20. J.H. LAU, HANDBOOK OF TAPE AUTOMATED BONDING, VAN NOSTRAND REINHOLD, 1992 Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

Soldering Process Parameters The parameters that affect wetting and spreading phenomena include: • • • • •

TEMPERATURE TIME VAPOR PRESSURE METALLURGICAL AND CHEMICAL NATURE OF THE SURFACES GEOMETRY OF THE SOLID

The manipulation of each parameter can result in some control of the wetting and soldering processes. Temperature is important to wetting and spreading for several reasons. Higher temperatures result in greater atomic

activity and provide some of the energy needed to overcome surface barriers. Higher temperatures also increase reaction rates exponentially. Temperature determines phase relationships for a given alloy composition, whether for the formation of a solid solution or for that of an intermetallic compound. The fluidity of the liquid solder is also increased with increasing temperature. In addition, higher temperatures promote the formation of oxides, which can be troublesome when soldering. Time. Wetting is time dependent because of the rate of kinetics of the solder-substrate interaction, the flux action, and the

heat conduction of the substrate. The soldering process is time sensitive, because time is needed to provide sufficient opportunity for the solder to wet, penetrate, or "wick up" into the various areas of the substrate that must be incorporated into the solder joint. It is clear that measures of the time to initiate wetting and of the time it takes for spreading to occur are needed. Vapor pressure is considered to be negligible for lead-tin solders, but it can be a significant factor when working with

certain solders. Caution must be exercised when selecting and working with alloys that have appreciable vapor pressures. The chemical and metallurgical nature of a surface affects wetting and spreading in several ways. Alloy content, as

well as the fluidity of the liquid, both affect wetting and, thus, the rate of spreading. Grain boundaries are wet differently than the bulk of the metal. The formation of solid solutions or intermetallic compounds also can have a considerable effect on wetting and spreading. The presence of oxide skins will prevent wetting and spreading, because the solder generally will not wet oxides. Likewise, films or particles of organic matter will interfere with wetting. Oils and silicones are frequent offenders and usually require different types of cleaning materials for their removal than the types used for oxides. Surface Geometry. Because contact angle is a critical measure of wetting and spreading behavior, surface geometry is

an important factor in the control of spreading. Soldering artisans have long used a wire brush not only to clean the surface to be soldered, but also to aid the spreading of the solder by providing sharp grooves that act as small capillaries to assist wetting. Isotropic spreading is actually what is needed. Another condition of surface interaction that is sometimes encountered in soldering is dewetting, where a metal surface is initially wet by solder, but as the solder cools, its cohesive forces exceed the forces involved in wetting, and the solder balls up on the metal surface. This usually results from surface contamination of the metal or from the entrapment of foreign particles in the metal surface, leading to weak wetting or, possibly, only local wetting. Base-Metal Selection. A sound soldered joint is achieved by selecting and using the proper materials and processes.

Base metals are usually selected to achieve the specific property requirements of a component. These properties can include strength, ductility, electrical conductivity, weight, corrosion resistance, and others. When soldering is required, the solderability of the base materials should also be a selection factor. Both flux selection and surface preparation will be affected by the solderability of the base materials to be joined.

The solderability of metals and alloys is not simply a matter of chemical nobility, as might be supposed when regarding the good solderability of the noble metals, which do not readily form oxide or tarnish films. Although cadmium and tin both form oxides readily, they are considered easy to solder. On the other hand, chromium, nickel, and aluminum also form oxide films readily, but are difficult to solder. The difference is in the extremely adherent, protective nature of the oxides formed on chromium, nickel, and aluminum, compared with the oxides that form on tin and cadmium (Table 2).

TABLE 2 RELATIVE SOLDERABILITY OF SELECTED METALS AND ALLOYS

EASY TO SOLDER • • • • • • •

PLATINUM GOLD COPPER SILVER CADMIUM PLATE TIN SOLDER PLATE

LESS EASY TO SOLDER • • • • • •

LEAD NICKEL PLATE BRASS BRONZE RHODIUM BERYLLIUM COPPER

DIFFICULT TO SOLDER • • • •

GLAVANIZED IRON TIN-NICKEL NICKEL-IRON MILD STEEL

VERY DIFFICULT TO SOLDER • • • •

CHROMIUM NICKEL-CHROMIUM NICKEL-COPPER STAINLESS STEEL

MOST DIFFICULT TO SOLDER • •

ALUMINUM ALUMINUM BRONZE

NOT SOLDERABLE • •

BERYLLIUM TITANIUM

It should be noted that chromium, nickel, and aluminum are all soldered regularly with good results, but that special attention must be given to the selection of fluxes, which must be very active. In many cases, the use of active fluxes is either restricted or not allowed. Therefore, these hard-to-solder metals and alloys always require special consideration in order to provide reproducible soldering. Numerous tests and methods to measure solderability are available. The tendencies of solder to either wet or spread on a given material are critical when evaluating candidate soldering systems. When hand soldering on single, large-dimension bodies prevailed, the skilled artisan could adjust parameters while watching the results develop. In most cases, visual determinations of wetting and spreading characteristics provided a sufficient basis for process control. Today, operating parameters must be controlled very carefully, because the soldering of components and/or electronic assemblies can involve many joints that are soldered at one time, or joints that are often hidden, or dimensions that are extremely small, or situations where joint-to-joint accommodations cannot be made. This necessitates the use of automatic soldering systems. Moreover, conditions must be uniform over the whole structure being soldered. In this environment, it is essential that the processes be based on measurements of wetting and spreading behavior, so that the controlling parameters can be accurately set. Additional information is available in the articles "Brazeability and Solderability of Engineered Materials," "General Soldering," and "Soldering in Electronic Applications" in this Volume. Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

Soldering Equipment The proper application of heat is of paramount importance in any soldering operation. The solder should melt while the surface is heated to permit the molten solder to wet and flow over the surface. The best heating method is therefore another important consideration. Soldering Irons. The traditional soldering tool is the soldering iron, or bit, with a copper tip than can be heated electrically, by direct flame, or in an oven. Because soldering is a heat-transfer process, the maximum surface area of the heated tip should contact the base metal. The solder itself should not be melted upon the tip of the iron when a joint is being made. To lengthen the usable life of a copper tip, a coating of solder-wettable metal, such as iron with or without additional coatings, is applied to the surface of the copper. The rate of dissolution of the iron coating in molten solder is substantially less than the rate for copper. The iron coating also shows less wear, oxidation, and pitting than uncoated copper.

Soldering irons are available in a large variety of sizes and designs, ranging from a small pencil to special irons or bits that weigh 2 kg (5 lb) or more. The selection of the iron depends on the application and the quantity of heat needed at the joint. The heat recovery time of the iron should be fast enough to keep up with the job. Regardless of the heating method, the tip performs the following functions: • • • • •

STORING AND CONDUCTING HEAT FROM THE HEAT SOURCE TO THE PARTS BEING SOLDERED STORING MOLTEN SOLDER CONVEYING MOLTEN SOLDER WITHDRAWING SURPLUS MOLTEN SOLDER BRINGING THE WORKPIECE JOINT AREA TO THE SOLDERING TEMPERATURE

The angle at which the copper tip is applied to the work is important in terms of delivering maximum heat. The flat side of the tip should be applied to the work to obtain the maximum area of contact. Flux-cored solders should not be melted on the soldering tip, because this destroys the effectiveness of the flux. The cored solder should be touched to the

soldering tip to initiate good heat transfer, and then the solder should be melted on the work parts to complete the solder joint (Ref 11). Additional information is available in the article "Iron Soldering" in this Volume. Torch Soldering (TS). The selection of a gas torch is controlled by the size, mass, and configuration of the assembly to be soldered. When fast soldering is necessary, a flame is frequently used. The flame temperature is controlled by the fuel mixture used. Fuel gas burned with oxygen gives the highest flame temperature possible with that gas. The highest flame temperatures are attained with acetylene. Lower temperatures are obtained with propane, butane, natural gas, and manufactured gas.

Multiple flame tips, or burners, often have shapes that are suitable to the work. They can be designed to operate on oxygen and fuel gas, compressed air and fuel gas, or Bunsen-type torches. When adjusting tips or torches, care should be taken to avoid adjustments that result in a "sooty" flame, because the carbon deposited on the work will prevent the flow of solder (Ref 11, 15). Detailed information is available in the article "Torch Soldering" in this Volume. Dip soldering (DS) utilizes a molten bath of solder to supply both the heat and solder required to produce the joints

(Fig. 9a). This method is useful and economical, because an entire unit comprising any number of joints can be soldered in one operation after proper cleaning and fluxing. Fixtures are required to contain the unit and maintain proper joint clearances during solder solidification. The soldering pot should be large enough so that, at a given rate of production, the units being dipped will not appreciably lower the temperature of the solder bath.

FIG. 9 SELECTED SOLDERING TECHNIQUES USED FOR HIGH-VOLUME PRODUCTION APPLICATIONS. (A) DIPPING ON A STATIC BATH. (B) WAVE SOLDERING SYSTEM. (C) CASCADE SOLDERING SYSTEM

Additional information is available in the article "Dip Soldering" in this Volume. Wave Soldering (WS). Electronic components are commonly soldered to PWBs by wave soldering (Fig. 9b). The solder is pumped out of a narrow slot to produce a wave or series of waves. This approach, which lends itself to automatic soldering, involves a continuously pumped crest of solder that peaks at the PWB. Another method is "drag," or planar, soldering. This method uses a bath, into which the board is automatically dipped with controlled conveyorization (Ref 21, 22).

In cascade soldering (Fig. 9c), the solder flows down a trough because of gravity and is returned by pump to the upper reservoir. Both of these wave solder systems are excellent in that a virtually oxide-free solder surface is presented to the part, and the flux and vapors are dislodged by the flow of the solder. Integrated wave soldering systems for printed wiring assemblies provide units that can apply the flux, dry and preheat the board, solder components, and clean the completed assembly. Some systems have special features that apply the flux by passing the assemblies through a wave, by spraying, by rolling, or by dipping. Several systems use oil mixed with the solder to aid in eliminating icicles and bridging between conductor paths. Another system features dual waves that flow in the opposite direction to the travel of the board. Additional information is available in the article "Wave Soldering" in this Volume. Induction Soldering (IS). The only requirement for a material that is to be induction soldered is that it be an electrical

conductor. The rate of heating depends on the induced current flow. The distribution of heat obtained with induction heating is a function of the induced current frequency. The higher frequencies concentrate the heat at the surface of the workpiece. Several types of equipment are available for induction heating: the vacuum tube oscillator, the resonant spark gap, the motor-generator unit, and solid-state units. Induction heating is generally applicable for soldering operations, with the following requirements and attributes: • • • • •

LARGE-SCALE PRODUCTION APPLICATION OF HEAT TO A LOCALIZED AREA MINIMUM OXIDATION OF SURFACE ADJACENT TO THE JOINT GOOD APPEARANCE AND CONSISTENTLY HIGH JOINT QUALITY SIMPLE JOINT DESIGN THAT LENDS ITSELF TO MECHANIZATION

The induction technique requires that the parts being joined have clean surfaces and that joint clearances be maintained accurately. High-grade solders are generally required to obtain rapid spreading and good capillary flow. Preforms often afford the best means of supplying the correct amount of solder and flux to the joint. When induction soldering dissimilar metals (particularly joints composed of both magnetic and nonmagnetic components), attention must be given to the design of the induction coil in order to bring both parts to approximately the same temperature. Additional information is available in the article "Induction Soldering" in this Volume. Furnace Soldering. There are many applications, especially in high-production soldering, where a furnace will produce

consistent and satisfactory soldering. Although this method is not widely used, furnace heating should be considered under the following circumstances:

• • •

WHEN ENTIRE ASSEMBLIES CAN BE BROUGHT TO THE SOLDERING TEMPERATURE WITHOUT DAMAGE TO ANY OF THE COMPONENTS WHEN PRODUCTION IS SUFFICIENTLY LARGE TO ALLOW EXPENDITURES FOR JIGS AND FIXTURES TO HOLD THE PARTS DURING THE SOLDERING PROCESS WHEN THE ASSEMBLY IS COMPLICATED IN NATURE, MAKING OTHER HEATING METHODS IMPRACTICAL

The proper clamping fixtures are very important during furnace soldering. Movement of the joint during solder solidification can result in a poor joint. Another important consideration in furnace soldering is flux selection. Rosin and organic fluxes are subject to decomposition when maintained at elevated temperatures for an extended period of time. When either rosin or organic flux is used, the part must be brought rapidly to the liquidus temperature of the solder. It is sometimes beneficial to dip the parts in a hot flux solution before placing them in the oven. When using rosin-base flux, it is generally necessary to use a solder with a tin content of 50% or more. The reducing atmosphere used in the furnace does not allow joints to be made without flux, because the temperatures at which these atmospheres become reducing are far above the liquidus temperature of the solders. The use of inert atmospheres will prevent further oxidation of the parts, but flux must be used to remove the oxide that is already present. The furnaces should be equipped with adequate temperature controls because solder flow begins at approximately 45 to 50 °C (80 to 90 °F) above the solder liquidus temperature. The optimum condition exists when the heating capacity of the furnace is sufficient to heat the parts rapidly to the liquidus temperature of the solder. Detailed information is available in the article "Furnace and Infrared Soldering" in this Volume. Resistance soldering (RS) involves placing the workpiece either between a ground and a movable electrode or

between two movable electrodes as part of an electrical circuit. Heat is applied to the joint both by the electrical resistance of the metal being soldered and by conduction from the electrode, which is usually carbon. Production assemblies can utilize multiple electrodes, rolling electrodes, or special electrodes, depending on which setup offers the most advantageous soldering speed, localized heating, and power consumption. Resistance soldering electrode bits generally cannot be tinned, and the solder must be fed directly into the joint. Additional information is available in the article "Resistance Soldering" in this Volume. Infrared Soldering (IRS). Optical soldering systems that focus infrared light (radiant energy) on the joint by means of

a lens are available. Lamps with power ratings that range from 45 to 1500 W (140 to 4700 Btu/h) can be used for different application requirements. The devices can be programmed through a silicon-controlled rectifier (SCR) power supply with an internal timer. The most common sources of infrared heating for soldering applications are heated filaments. The quartz-iodine tungstenfilament lamp is widely used because it is very stable and reliable over a wide range of temperatures. In general, infrared soldering systems are simple and inexpensive to operate. One of the most critical operating parameters is surface condition. Variations in the condition of the solder surface can be compensated for, to some extent, by adjusting the heating power. Advantages are process repeatability, ability to concentrate or focus the energy with reflectors and lenses, economy of operation, and absence of contact with the workpiece. Detailed information is available in the article "Furnace and Infrared Soldering" in this Volume. Ultrasonic Soldering. This process uses a transducer as the source of ultrasonic energy. The transducer is energized in a bath of molten solder, and sound waves are coupled between the transducer and the workpiece, allowing the oxides in the base metal to be disrupted so that the solder melts the base metal. Sound waves are transmitted throughout the base metal, permitting wetting to occur on surfaces that are "blind" to the source. Ultrasonic soldering is also used to apply solderable coatings on difficult-to-solder metals.

Additional Soldering Processes. In addition to the above-mentioned processes, hot-gas soldering, spray-gun soldering,

and condensation soldering are also used. Hot gas soldering utilizes a fine jet of inert gas, heated to above the liquidus of the solder. The gas acts as a heat-

transfer medium and as a blanket to reduce access of air at the joint. Additional information is available in the article "Hot Gas Soldering" in this Volume. Spray gun soldering is a heating method used when the contour of the part to be soldered is difficult to follow with

either a wiping or drop method or when the part is placed in the assembly in such a way that the solder cannot be applied after the parts are assembled. Gas-fired or electrically heated guns are available. Each type is designed to spray molten or semimolten solder on the work from a continuously fed solid solder wire. Soldering guns use either propane mixed with oxygen or natural gas mixed with air to heat and to spray a continuously fed solid solder wire of approximately 3.2 mm (0.12 in.) diameter. About 90% of the solder wire is melted by the flame of the gun. The solder contacts the workpiece in a semiliquid form. The workpiece then supplies the balance of the heat required to melt and flow the solder. Adjustments can be made within the spray gun to control the solder spray. Vapor-phase soldering (also known as condensation soldering) utilizes the latent heat of vaporization of a condensing

saturated liquid to provide the heat required for soldering. A reservoir of saturated vapor over a boiling liquid provides a constant controlled temperature with rapid heat transfer. This method is useful for large assemblies, as well as for temperature-sensitive parts. Detailed information is available in the article "Vapor-Phase Soldering" in this Volume.


11. R.W. WOODGATE, THE HANDBOOK OF MACHINE SOLDERING, 2ND ED., JOHN WILEY & SONS, 1988 15. E. LIEBERMAN, MODERN SOLDERING AND BRAZING TECHNIQUES, BUSINESS NEWS PUBLISHING CO., 1988 21. P.J. BUD, MASS PRODUCTION TECHNIQUES USING THE PRINCIPLE OF WAVE SOLDERING, WELD. J., VOL 52 (NO. 7), JULY 1973, P 431-439 22. P.J. BUD AND D.A. ELLIOT, WAVESOLDERING AT ELEVATED TEMPERATURES, WELD. J., VOL 53 (NO. 2), FEB 1974, P 79-87 Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

Quality Control Table 3 lists that are commonly used to evaluate the solderability properties of selected soldered components.

TABLE 3 TEST STANDARDS USED TO EVALUATE SOLDERABILITY

BASIC METHOD DIP AND LOOK

TEST STANDARD ANSI/IPC-S-804 (BOARDS) ANSI/IPC-S-805 (COMPONENTS)

IEC 68-2-20 MIL-STD-202, 750, 883 EIA RS-186-9E ROTARY DIP ANSI-IPC-S-804 (BOARDS) IEC 68-2-20 BS4025 (BRITISH STANDARD) WETTING BALANCE ANSI-IPC-S-805 (COMPONENTS) IEC-68-2-20 MIL-STD-883 GLOBULE TEST ANSI-IPC-S-805 (COMPONENTS) IEC 68-2-20 BS 2011 (BRITISH STANDARD) DIN 40046 (GERMAN STANDARD) MENISCUS RISE NO STANDARDS TO DATE TIMED SOLDER RISE ANSI/IPC-S-804 (BOARDS) Dip Test. The most useful test for assessing solderability is the dip test, because of its accuracy and simplicity. It is

conducted with the additions of cams and timers so that standardized conditions are established. The dip test is a reasonable simulation of practical soldering conditions and starts with a test specimen that has been dipped in a mild flux, as is usually the case with soldering tests. The dip test involves immersing the test specimen into a clean solder bath, waiting an appropriate length of time, and then withdrawing the specimen from the bath. In most cases, the test is automated to ensure reproducibility. The results are evaluated by means of visual inspection and comparison with a set of standards. Thus, the dip test is operator sensitive, somewhat subjective, and of limited value, because it is not quantitative. Rotary Dip Test. This test was devised to provide a better simulation of the dynamic conditions that exist in machine

soldering operations. It is used to evaluate the solderability of PWBs, and involves bringing the specimen into moving contact with the molten solder for a specified period of time. This is a "go/no-go" test that requires complete wetting of the specimen surface within a set time, and is thus dependent on visual inspection of the test specimen. Wetting Balance (Meniscograph). Another common test employs the wetting balance, which gives a numerical result.

The equipment consists of a balance that records the forces exerted on a fluxed test specimen during a controlled dipping cycle in a clean solder bath. From the values obtained, it is possible to compare the rate of wetting, as well as the forces involved for various flux formulations. The speed of wetting is an especially important factor in the wave soldering process, where automation requires a timebased response of flux and solder as the work moves through the process machine at a constant speed. The wetting balance test cannot be used on finished, etched printed circuits or on single-sided circuits. It is often used on specimens cut from double-sided clad materials prior to etching. Although this test gives numerical results, they are not strictly quantitative. However, the results do provide the basis for a meaningful comparison among fluxes. Additional Tests. There are several other tests that provide practical information for special purposes, such as the

spreading test and the globule test. These tests are sometimes used to evaluate incoming materials or to aid in the selection of materials. They are not universally employed. The spreading test is used to provide a numerical rating for both fluxes and solders. It is described in detail in MIL-F-

14256, "Flux, Soldering, Liquid (Rosin Base)," and in QQ-S-571, "Solder, Electronic," which are national-level specifications. The globule test was once widely applied, but is now limited to being used to determine the wetting of component leads

and wires (Ref 24). Additional information is available in the article "Evaluation and Quality Control of Soldered Joints" in this Volume.


24. IEC-68-2-20, INTERNATIONAL ELECTROCHEMICAL COMMISSION, GENEVA Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

Future Outlook The quality soldered joints obtained by automation have resulted in the widespread use of automated equipment throughout the electronics industry, as well as throughout the world (Ref 25). When compared to processes such as adhesive bonding, welding, brazing, or mechanical joining, soldering offers many advantages. Automated soldering equipment for wave or planar soldering can be easily and economically installed. The high reliability that can be achieved with soldering results in a quality end product. Soldering offers the additional benefit of providing sequential assembly, a common practice in the electronics industry.


25. REFLOW SOLDERING, WELD. J., VOL 52 (NO. 1), JAN 1973, P 22-30 Fundamentals of Soldering Mel M. Schwartz, Sikorsky Aircraft

References

1. J. WOLTERS, ZUR GESEHICHTE DER LOTTECHNIK, DEGUSSA, WEST GERMANY, 1977 2. H.H. MANKO, SOLDERS AND SOLDERING, 2ND ED., MCGRAW-HILL, 1964 3. R.J. KLEIN WASSINK, SOLDERING IN ELECTRONICS, ELECTROCHEMICAL PUBLICATIONS LTD., 1989 4. L.P. LAMBERT, SOLDERING FOR ELECTRONIC ASSEMBLIES, MARCEL DEKKER, 1987 5. SOLDERING MANUAL, AWS, 1978 6. A.D. ROMIG, JR., Y.A. CHANG, J.J. STEPHENS, ET AL., PHYSICAL METALLURGY OF SOLDERSUBSTRATE REACTIONS, SOLDER MECHANICS, THE MINERALS, METALS, & MATERIALS SOCIETY, 1991, P 30-32 7. "SOFT SOLDER ALLOYS--CHEMICAL COMPOSITIONS AND FORMS," ISO 9453, 1ST ED., 1990 8. "SOFT SOLDERING FLUXES--CLASSIFICATION AND REQUIREMENTS--PART I: CLASSIFICATION, LABELING AND PACKAGING," ISO 9454-1, 1990 9. J.F. SHIPLEY, INFLUENCE OF FLUX, SUBSTRATE AND SOLDER COMPOSITION ON SOLDER WETTING, WELD. J., VOL 54 (NO. 10), OCT 1975, P 357S-362S 10. H.H. MANKO, SOLDERING FLUXES--PAST AND PRESENT, WELD. J., VOL 52 (NO. 3), MARCH 1973, P 163-166 11. R.W. WOODGATE, THE HANDBOOK OF MACHINE SOLDERING, 2ND ED., JOHN WILEY & SONS, 1988

12. T. THOMPSON, CONDENSATION/VAPOR-PHASE REFLOW SOLDERING, ASS. ENG., JUNE 1977, P 44-47 13. T.Y. CHU, A GENERAL REVIEW OF MASS SOLDERING METHODS, INS. CIRC., NOV 1976 14. R.C. PFAHL, JR., J.C. MOLLENDORF, AND T.Y. CHU, CONDENSATION SOLDERING, WELD. J., VOL 54 (NO. 1), 1975 15. E. LIEBERMAN, MODERN SOLDERING AND BRAZING TECHNIQUES, BUSINESS NEWS PUBLISHING CO., 1988 16. "SOLDER METAL SPECIFICATION," ASTM B 32, ANNUAL BOOK OF ASTM STANDARDS, PART B, ASTM 17. D.R. FREAR, W.B. JONES, AND K.R. KINSMAN, ED., SOLDER MECHANICS, TMS, 1991, P 30-41 18. PROC. IEEE 41ST ELECT. COMP. & TECH. CONF., IEEE, 1991 19. J. BROUS, EVALUATION OF POST-SOLDER FLUX REMOVAL, WELD. J., VOL 54 (NO. 12), 1975, P 444S-448S 20. J.H. LAU, HANDBOOK OF TAPE AUTOMATED BONDING, VAN NOSTRAND REINHOLD, 1992 21. P.J. BUD, MASS PRODUCTION TECHNIQUES USING THE PRINCIPLE OF WAVE SOLDERING, WELD. J., VOL 52 (NO. 7), JULY 1973, P 431-439 22. P.J. BUD AND D.A. ELLIOT, WAVESOLDERING AT ELEVATED TEMPERATURES, WELD. J., VOL 53 (NO. 2), FEB 1974, P 79-87 23. ELECTRONIC MATERIALS HANDBOOK, VOL 1, PACKAGING, ASM INTERNATIONAL, 1989 24. IEC-68-2-20, INTERNATIONAL ELECTROCHEMICAL COMMISSION, GENEVA 25. REFLOW SOLDERING, WELD. J., VOL 52 (NO. 1), JAN 1973, P 22-30 Introduction to Solid-State Welding Ray Dixon, Los Alamos National Laboratory

Solid-State Welding (SSW) processes are those that produce coalescence of the faying surfaces at temperatures below the melting point of the base metal being joined without the addition of brazing or solder filler metal. Pressure may or may not be applied. These processes involve either the use of deformation or of diffusion and limited deformation in order to produce high-quality joints between both similar and dissimilar materials. Dissimilar metal joints are necessary in applications that require a variety of material properties within the same component. For example, heat exchangers often require different types of stainless steels at each end, because of temperature-induced corrosion. Under laboratory conditions, dissimilar materials can be chosen based on physical or material properties that influence the phenomenon being studied. For whatever reason, and appropriate method of producing dissimilar metal joints can usually be determined (assuming it is even possible) by examining the phase diagram. If the diagram indicates difficulty in joining the materials (intermetallics, and so on), then a solid-state (nonmelting) process may be applicable. When a nonmelting process is chosen, it is only successful if a relatively strong joint is produced. There is considerable interest is quantifying joint strength, and extensive literature (Ref 1, 2, 3, 4, 5, 6, 7, 8, 9, 10, 11, 12, 13, 14) exists to help design and evaluate various specimen geometries. This experience includes not only techniques for conducting the tests, but analysis and interpretation of the data with respect to the physical properties of individual materials. The study of dissimilar welds necessarily leads to studies of the interfacial region, because it is the most likely site of part failures. In a two-component system (materials A and B), interfacial failure can be a result of the individual properties of either A and B, but they are by necessity synergistic. These synergisms and the resulting complex interactions that they produce are what is significant. From an engineering and processing standpoint, an interfacial "cookbook" that provides

all the answers would be ideal. However, this capability is not yet available, and, as explained in this Section, researchers are still far from reaching that goal. This is the case even for simple systems, such as metal pair that form solid solutions. The opportunities for understanding interfacial behavior come from studies conducted at the atomic level (for example, using scanning tunneling microscopy, transmission electron microscopy, first principles energy band calculations, and embedded atom potential calculations) with appropriate correlations to both micromechanics and a macroscopic level, and vice versa. The computational and experimental tools that begin to provide insight at these levels have only recently been developed. Because the interest in interfaces is so diverse (all combinations of metals, ceramics, semiconductors, glasses, and conducting polymers), the advances in interfacial science encompass the creative efforts of a large number of people and organizations. Thus, the quantity of available literature seems overwhelming, and perhaps, disjointed. It will become apparent that there is no preferred approach to the treatment of dissimilar welds or welding. It really depends on the interests and objectives of each engineer. For example, if you know that what you are doing actually works (that is, produces a joint) and are searching for understanding, then this Section of the Handbook will guide you to the appropriate technical literature. If you have an idea and need guidance for bonding, testing, processing, or interpreting the data, then this Section should also provide that direction. Lastly, if the reader wants state-of-the-art research and development topics, then those can be found in this Section, as well. This should provide the reader with a comprehensive overview of the welding of dissimilar metals. Therefore, this Section is concerned with the fundamentals of welding and joining materials via the application of a nonmelting process. The specific processes usually associated with this technology are also discussed. Before considering nonmelting processes in detail, however, a brief discussion of adhesion is presented. This is followed by articles that discuss the fundamentals of the major solid-state welding techniques.

References

1. P.G. CHARALAMBIDES, J. LUND, A.G. EVANS, AND R.M. MCMEEKING, "A TEST FOR DETERMINING THE FRACTURE RESISTANCE OF BIMATERIAL INTERFACES," REPORT M87-4, MATERIALS DEPARTMENT, COLLEGE OF ENGINEERING, UC SANTA BARBARA 2. Z. SUO AND J.W. HUTCHINSON, "INTERFACE CRACK BETWEEN TWO ELASTIC LAYERS," REPORT MECH-118, DIVISION OF APPLIED SCIENCES, HARVARD UNIVERSITY, MARCH 1988 3. N.P. O'DOWD, C.F. SHIH, AND M.G. STOUT, TEST GEOMETRIES FOR MEASURING INTERFACIAL FRACTURE TOUGHNESS, FOR INT. J. SOLIDS STRUCT., VOL 29, 1992, P 571-589 4. J.W. HUTCHINSON AND Z. SUO, MIXED MODE CRACKING IN LAYERED MATERIALS, ADV. IN APPL. MECH., VOL 28, J.W. HUTCHINSON AND T.Y. WU, ED., ACADEMIC PRESS, 1990 5. H.M. JENSEN, MIXED MODE INTERFACE FRACTURE CRITERIA, ACTA METALL. MATER., VOL 38 (NO. 12), 1990, P 2637-2644 6. R.L. BRADY, R.S. PORTER, AND J.A. DONOVAN, A BUCKLED PLATE TEST TO MEASURE INTERFACIAL TOUGHNESS IN COMPOSITES, J. MATER. SCI., VOL 24, 1989, P 4138-4143 7. M.-Y. HE AND J.W. HUTCHINSON, "KINKING OF A CRACK OUT OF AN INTERFACE," REPORT MECH-113, DIVISION OF APPLIED SCIENCE, HARVARD UNIVERSITY, FEB 1988 8. J.J. MECHOLSKY AND L.M. BAKER, A CHEVRON-NOTCHED SPECIMEN FOR FRACTURE TOUGHNESS MEASUREMENTS OF CERAMIC-METAL INTERFACES, CHEVRON-NOTCHED SPECIMENS: TESTING AND STRESS ANALYSIS, STP 855, J.H. UNDERWOOD, S.W. FREIMAN, AND F.I. BARATTA, ED., ASTM, 1984, P 324-336 9. W.W. GERBERICH, D.L. DAVIDSON, AND M. KACZOROWSKI, EXPERIMENTAL AND THEORETICAL STRAIN DISTRIBUTIONS FOR STATIONARY AND GROWING CRACKS, J. MECH. PHYS. SOLIDS, VOL 38 (NO. 1), 1990, P 87-113 10. J. QU AND Q. LI, INTERFACIAL DISLOCATION AND ITS APPLICATIONS TO INTERFACE CRACKS IN ANISOTROPIC BIMATERIALS, J. ELAST., VOL 26, 1991, P 169-195

11. A.N. STROH, DISLOCATIONS AND CRACKS IN ANISOTROPIC ELASTICITY, PHILOS. MAG., VOL 7, 1958, P 625-646 12. Q.Q. LI AND T.C.T. TING, LINE INCLUSION IN ANISOTROPIC ELASTIC SOLIDS, J. APPL. MECH., VOL 56, 1989, P 556-563 13. T.C.T. TING, EXPLICIT SOLUTION AND INVARIANCE OF THE SINGULARITIES AT AN INTERFACE CRACK IN ANISOTROPIC COMPOSITES, INT. J. SOLIDS STRUCT., VOL 22, 1986, P 965-983 14. M.D. PASHLEY AND D. TABOR, ADHESION AND DEFORMATION PROPERTIES OF CLEAN AND CHARACTERIZED METAL MICRO-CONTACTS, VACUUM, VOL 10-12, 1981, P 619-623 Introduction to Solid-State Welding Ray Dixon, Los Alamos National Laboratory

References

1. P.G. CHARALAMBIDES, J. LUND, A.G. EVANS, AND R.M. MCMEEKING, "A TEST FOR DETERMINING THE FRACTURE RESISTANCE OF BIMATERIAL INTERFACES," REPORT M87-4, MATERIALS DEPARTMENT, COLLEGE OF ENGINEERING, UC SANTA BARBARA 2. Z. SUO AND J.W. HUTCHINSON, "INTERFACE CRACK BETWEEN TWO ELASTIC LAYERS," REPORT MECH-118, DIVISION OF APPLIED SCIENCES, HARVARD UNIVERSITY, MARCH 1988 3. N.P. O'DOWD, C.F. SHIH, AND M.G. STOUT, TEST GEOMETRIES FOR MEASURING INTERFACIAL FRACTURE TOUGHNESS, FOR INT. J. SOLIDS STRUCT., VOL 29, 1992, P 571-589 4. J.W. HUTCHINSON AND Z. SUO, MIXED MODE CRACKING IN LAYERED MATERIALS, ADV. IN APPL. MECH., VOL 28, J.W. HUTCHINSON AND T.Y. WU, ED., ACADEMIC PRESS, 1990 5. H.M. JENSEN, MIXED MODE INTERFACE FRACTURE CRITERIA, ACTA METALL. MATER., VOL 38 (NO. 12), 1990, P 2637-2644 6. R.L. BRADY, R.S. PORTER, AND J.A. DONOVAN, A BUCKLED PLATE TEST TO MEASURE INTERFACIAL TOUGHNESS IN COMPOSITES, J. MATER. SCI., VOL 24, 1989, P 4138-4143 7. M.-Y. HE AND J.W. HUTCHINSON, "KINKING OF A CRACK OUT OF AN INTERFACE," REPORT MECH-113, DIVISION OF APPLIED SCIENCE, HARVARD UNIVERSITY, FEB 1988 8. J.J. MECHOLSKY AND L.M. BAKER, A CHEVRON-NOTCHED SPECIMEN FOR FRACTURE TOUGHNESS MEASUREMENTS OF CERAMIC-METAL INTERFACES, CHEVRON-NOTCHED SPECIMENS: TESTING AND STRESS ANALYSIS, STP 855, J.H. UNDERWOOD, S.W. FREIMAN, AND F.I. BARATTA, ED., ASTM, 1984, P 324-336 9. W.W. GERBERICH, D.L. DAVIDSON, AND M. KACZOROWSKI, EXPERIMENTAL AND THEORETICAL STRAIN DISTRIBUTIONS FOR STATIONARY AND GROWING CRACKS, J. MECH. PHYS. SOLIDS, VOL 38 (NO. 1), 1990, P 87-113 10. J. QU AND Q. LI, INTERFACIAL DISLOCATION AND ITS APPLICATIONS TO INTERFACE CRACKS IN ANISOTROPIC BIMATERIALS, J. ELAST., VOL 26, 1991, P 169-195 11. A.N. STROH, DISLOCATIONS AND CRACKS IN ANISOTROPIC ELASTICITY, PHILOS. MAG., VOL 7, 1958, P 625-646 12. Q.Q. LI AND T.C.T. TING, LINE INCLUSION IN ANISOTROPIC ELASTIC SOLIDS, J. APPL. MECH., VOL 56, 1989, P 556-563 13. T.C.T. TING, EXPLICIT SOLUTION AND INVARIANCE OF THE SINGULARITIES AT AN INTERFACE CRACK IN ANISOTROPIC COMPOSITES, INT. J. SOLIDS STRUCT., VOL 22, 1986, P 965-983 14. M.D. PASHLEY AND D. TABOR, ADHESION AND DEFORMATION PROPERTIES OF CLEAN AND CHARACTERIZED METAL MICRO-CONTACTS, VACUUM, VOL 10-12, 1981, P 619-623

Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Introduction IN ANY SOLID-STATE (NONMELTING) WELDING PROCESS, there are two primary areas of concern: Will the material bond and how strong is the bond? Literature that deals with the science and technology of material compatibility and bond strength is readily available. This article will review recent work in quantifying adhesion, bonding, and interfacial characterization and strength. Subsequent articles in this Section of the Handbook focus on the engineering principles associated with the formation of welds using these techniques. This discussion concentrates on metal-metal configurations, with only passing comments on metal-ceramic and metal/ceramic-semiconductor interfaces. This restriction is necessary to limit the length of this article. Wherever appropriate, applications of thin films on substrates will also be used. Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Adhesion Energy Surface energy, Γsurf, is defined as the increase in energy, ∆E, of a system divided by the area, A, caused by the creation of the surface:

ΓSURF = Σ∆EI/A

(EQ 1)

The atoms that contribute to the extra energy, ∆Ei, are usually within 10 lattice parameters or less of the surface (or interface) (Ref 1, 2, 3). A freshly created surface has a high surface energy, whereas a relaxed or reconstructed one is lower, by about 3 to 5%. At 0 K (-459 °F), surface energies vary about 15 to 25% for planes of various indices. The variation is less at higher temperatures, because of entropy. For example, the (111) surface of a face-centered cubic (fcc) structure has the lowest energy, whereas the (210) surface is about 15% higher (Ref 1, 2, 3). The ideal fracture energy, coh, which is used in estimating the Griffith criterion, is twice the surface energy:

ΓCOH (HKL) = 2 · ΓSURF (HKL)

(EQ 2)

where (hkl) are the indices of the plane. The actual energy consumed in fracture or crack propagation is composed of two terms: one is the Γcoh and the other is the plastic work consumed per unit area Γplastic (Ref 4, 5, 6, 7, 8). For ionic crystals, Γplastic is usually a few times Γcoh; for metals, Γcoh is about 1 to 4 J/m2 (2 × 10-5 to 1 × 10-4 cal/cm2); and Γplastic is about 1000 J/m2 (2 × 10-2 cal/cm2). For ceramic systems, Γcoh is about the same, but Γplastic is about 10 to 100 J/m2 (2 × 10-4 to 2 × 10-3 cal/cm2) (Ref 7). Γplastic and Γcoh are related, because when Γcoh is zero, no plastic work (Γplastic) will be consumed. The exact form of this dependence is not known, although some attempts have been made (Ref 4, 5, 6) to correlate the two terms.


1. S.P. CHEN, D.J. SROLOVITZ, AND A.F. VOTER, COMPUTER SIMULATION ON SURFACES AND [001] TILT GRAIN BOUNDARIES IN NI, AL, AND NI3AL, J. MATER. RES., VOL 4, 1989, P 62-77 2. S.P. CHEN AND A.F. VOTER, RECONSTRUCTION OF THE (310), (210), AND (110) SURFACES IN FCC METALS, SURF. SCI. LETT., VOL 244, 1991, P L107-L112 3. S.P. CHEN, THEORETICAL STUDIES OF METALLIC INTERFACES, MATER. SCI. ENG. B, VOL 6, 1990, P 113-121

4. M.L. JOKL, V. VITEK, AND C.J. MCMAHON, JR., A MICROSCOPIC THEORY OF BRITTLE FRACTURE IN DEFORMABLE SOLIDS: A RELATION BETWEEN IDEAL WORK TO FRACTURE AND PLASTIC WORK, ACTA METALL., VOL 28, 1980, P 1479-1488 5. M.L. JOKL, V. VITEK, AND C.J. MCMAHON, JR., ON THE MICROMECHANICS OF BRITTLE FRACTURE: EXISTING VS INJECTED CRACKS, ACTA METALL., VOL 37, 1989, P 87-97 6. J. HACK, S.P. CHEN, AND D. SROLOVITZ, A KINETIC CRITERION FOR QUASI-BRITTLE FRACTURE, ACTA METALL., VOL 37, 1989, P 1957-1970 7. A. KELLY AND N.H. MACMILLAN, STRONG SOLIDS, 3RD ED., OXFORD SCIENCE PUBLICATIONS, 1986 8. V. VITEK, MICROMECHANISMS OF INTERGRANULAR BRITTLE FRACTURE IN INTERMETALLIC COMPOUND, J. PHYS. III, VOL 1, 1991, P 1085-1097 Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Grain Boundary Energy The grain boundary energy, Γgb, is defined the same way as surface energy is, but it depends on the relative orientation of the adjacent grains (Ref 1):

ΓGB = Σ∆EI/AI (GB)

(EQ 3)

At misorientation angles of greater than 25°, Γgb usually reaches a plateau that characterizes random grain boundaries. Special grain boundaries, like Σ5 (210) or (310), tend to have slightly lower energies (Ref 1, 2, 3, 9). Twin boundary energies are even lower than special grain boundaries. For metals, Γgb usually ranges from 40 to 60% of the surface energy (Ref 10), whereas the twin energy is only about 5 to 15%. Grain boundary cohesive energy is defined as (Ref 1):

ΓCOH (GB) = ΓSURF (GRAIN A) + ΓSURF (GRAIN B) - ΓGB

(EQ 4)

For general and random grain boundaries, grain boundary cohesive energy is smaller. General grain boundaries tend to be weaker than the small angle boundaries.


1. S.P. CHEN, D.J. SROLOVITZ, AND A.F. VOTER, COMPUTER SIMULATION ON SURFACES AND [001] TILT GRAIN BOUNDARIES IN NI, AL, AND NI3AL, J. MATER. RES., VOL 4, 1989, P 62-77 2. S.P. CHEN AND A.F. VOTER, RECONSTRUCTION OF THE (310), (210), AND (110) SURFACES IN FCC METALS, SURF. SCI. LETT., VOL 244, 1991, P L107-L112 3. S.P. CHEN, THEORETICAL STUDIES OF METALLIC INTERFACES, MATER. SCI. ENG. B, VOL 6, 1990, P 113-121 9. S.P. CHEN, STUDIES OF IRIDIUM SURFACES AND GRAIN BOUNDARIES, PHIL. MAG. A, VOL 65, 1992, TO BE PUBLISHED 10. L.E. MURR, INTERFACIAL PHENOMENA IN METALS AND ALLOYS, ADDISON-WESLEY, 1975 Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Interfacial Energy By definition, the interfacial energy (Ref 2, 3) between two bulk materials (A and B) is the same as the grain boundary energy, except that there are two reference states:

ΓINTERFACE = [Σ∆EI (A) + Σ∆EJ(B)]/A

(EQ 5)

where ∆Ei(A) · [(∆Ej(B)] is the extra energy of atom i of A (B) type, compared to its bulk state of A (B) materials. The interfacial energy between similar metals tends to be small. For dissimilar materials, however, interfacial energies tend toward large positive values if they do not form compounds and large negative values if they do form compounds. The ideal adhesion energy is similar to grain boundary cohesion and is the difference between the two surfaces and the interfacial energies:

ΓADHESION = ΓSURFACE (A) + ΓSURFACE (B) - ΓINTERFACE (A/B)

(EQ 6)


2. S.P. CHEN AND A.F. VOTER, RECONSTRUCTION OF THE (310), (210), AND (110) SURFACES IN FCC METALS, SURF. SCI. LETT., VOL 244, 1991, P L107-L112 3. S.P. CHEN, THEORETICAL STUDIES OF METALLIC INTERFACES, MATER. SCI. ENG. B, VOL 6, 1990, P 113-121

Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Theory of Adhesion Rose et al. (Ref 11) have calculated the bimaterial interfacial cohesion for many combinations of elements using band structure methods. They found a universal scaling relationship for the adhesion of different materials that extends beyond metal-metal interfaces to nonmetal systems. A few calculations are necessary to identify the adhesion behavior of the system and set the appropriate energy, length, and elastic constant scales. Plastic contributions are not treated in this formalism. The plastic work is more likely to be solved by incorporating the atomistic simulation and continuum analysis of Jokl et al. (Ref 4, 5) and Hack et al. (Ref 6). The tight binding method give qualitative trends as the energy band is filled for calculations using the elements in the periodic table (Ref 12, 13, 14). The correct trend seems to be captured by the band-filling argument pioneered in the original work of Friedel (Ref 15). For noble metals, a rigid band picture is preferred (Ref 14). Another useful tool is the Pettifor structure map. By using the Mendelev number for each of the 103 elements (Ref 13, 16), Pettifor found that alloying tendencies can be grouped into separate regions. If the interfacial chemistry and average Mendelev number are known, then the interfacial structure can be found from the Pettifor map. Extensive atomistic simulation work, using effective medium theory (Ref 17, 18), embedded atom potentials (Ref 19), or local volume potentials (Ref 1, 2, 3, 8, 9, 20, 21, 22, 23) had been very useful in describing the properties of many metals and compounds and their interactions with impurities such as boron and sulfur in Ni3Al and NiAl (Ref 20, 21). These atomic interaction descriptions include pair potential and many body interactions and have been obtained for systems that include Ni3Al, Cu3Au, NiAl, nickel, aluminum, iridium, gold, platinum, silver, copper, molybdenum, tantalum, iron, tungsten, niobium, beryllium, titanium, hafnium, zirconium, and rhenium (Ref 1, 2, 3, 9, 20, 21, 22, 23, 24, 25, 26, 27). These studies indicate that stoichiometry (Ref 1, 20, 21) and impurities (Ref 20, 21) can be crucial in determining ideal and realistic adhesion, including effects induced by boron and sulfur. They represent the fundamentals for future work,

which is expected to have significant impact on the theory using atomistic simulation under well-controlled situations for all conditions of stress, temperature, and chemical environments. These calculations should enable direct linking of theoretical studies and materials by design concepts, as applied in NiAl microalloyed with boron or iron (Ref 24, 25, 26, 27). DeBoer et al. (Ref 28) have developed excellent tools for screening bonding materials, based on two properties of the elements: electronegativity and electron density. One can use the Miedema formula for the adhesion of a bimetallic interface, based on the work function (closely related to the Pauling electronegativity), Φ , and the electron density, n, at the Wigner-Seitz radius, as described in Ref 28.


1. S.P. CHEN, D.J. SROLOVITZ, AND A.F. VOTER, COMPUTER SIMULATION ON SURFACES AND [001] TILT GRAIN BOUNDARIES IN NI, AL, AND NI3AL, J. MATER. RES., VOL 4, 1989, P 62-77 2. S.P. CHEN AND A.F. VOTER, RECONSTRUCTION OF THE (310), (210), AND (110) SURFACES IN FCC METALS, SURF. SCI. LETT., VOL 244, 1991, P L107-L112 3. S.P. CHEN, THEORETICAL STUDIES OF METALLIC INTERFACES, MATER. SCI. ENG. B, VOL 6, 1990, P 113-121 4. M.L. JOKL, V. VITEK, AND C.J. MCMAHON, JR., A MICROSCOPIC THEORY OF BRITTLE FRACTURE IN DEFORMABLE SOLIDS: A RELATION BETWEEN IDEAL WORK TO FRACTURE AND PLASTIC WORK, ACTA METALL., VOL 28, 1980, P 1479-1488 5. M.L. JOKL, V. VITEK, AND C.J. MCMAHON, JR., ON THE MICROMECHANICS OF BRITTLE FRACTURE: EXISTING VS INJECTED CRACKS, ACTA METALL., VOL 37, 1989, P 87-97 6. J. HACK, S.P. CHEN, AND D. SROLOVITZ, A KINETIC CRITERION FOR QUASI-BRITTLE FRACTURE, ACTA METALL., VOL 37, 1989, P 1957-1970 8. V. VITEK, MICROMECHANISMS OF INTERGRANULAR BRITTLE FRACTURE IN INTERMETALLIC COMPOUND, J. PHYS. III, VOL 1, 1991, P 1085-1097 9. S.P. CHEN, STUDIES OF IRIDIUM SURFACES AND GRAIN BOUNDARIES, PHIL. MAG. A, VOL 65, 1992, TO BE PUBLISHED 11. J.H. ROSE, J.R. SMITH, F. GUINEA, AND J. FERRANTE, UNIVERSAL FEATURES OF THE EQUATION OF STATE OF METALS, PHYS. REV. B, VOL 29, 1984, P 2963-2969 12. T.B. MASSALSKI, ED., BINARY ALLOYS PHASE DIAGRAMS, ASM, 1986, P 2176-2181 13. D.G. PETTIFOR, ELECTRON THEORY OF METALS, PHYSICAL METALLURGY, R.W. CAHN AND P. HAASEN, ED., NORTH-HOLLAND, AMSTERDAM, 1983, P 73-152 14. J.S. FAULKNER, THE MODERN THEORY OF ALLOYS, PROG. MATER. SCI., VOL 27, 1982, P 1-187 15. J. FRIEDEL, TRANSITION METALS, ELECTRONIC STRUCTURE OF THE D-BAND, ITS ROLE IN THE CRYSTALLINE AND MAGNETIC STRUCTURES, THE PHYSICS OF METALS, J.M. ZIMAN, ED., CAMBRIDGE UNIVERSITY PRESS, 1969, P 340-408 16. D.G. PETTIFOR, NEW MANY-BODY POTENTIAL FOR THE BOND ORDER, PHYS. REV. LETT., VOL 63, 1989, P 2480-2483 17. J.K. NORSKOV AND N.D. LANG, EFFECTIVE-MEDIUM THEORY OF CHEMICAL BINDING: APPLICATION TO CHEMISORPTION, PHYS. REV. B, VOL 21, 1980, P 2131-2136 18. M.J. STOTT AND E. ZAREMBA, QUASIATOMS: AN APPROACH TO ATOMS IN NONUNIFORM ELECTRONIC SYSTEMS, PHYS. REV. B, VOL 22, 1980, P 1564-1583 19. M.S. DAW AND M.I. BASKES, EMBEDDED-ATOM METHOD: DERIVATION AND APPLICATION TO IMPURITIES, SURFACES, AND OTHER DEFECTS IN METALS, PHYS. REV. B, VOL 29, 1984, P 6443-6453 20. S.P. CHEN, A.F. VOTER, R.C. ALBERS, A.M. BORING, AND P.J. HAY, INVESTIGATION OF BORON EFFECT ON NI AND NI3AL GRAIN BOUNDARIES, J. MATER. RES., VOL 5, 1990, P 955970

21. S.P. CHEN, A.F. VOTER, R.C. ALBERS, A.M. BORING, AND P.J. HAY, THEORETICAL STUDIES OF GRAIN BOUNDARIES IN NI3AL WITH BORON OR SULFUR, SCR. METALL., VOL 23, 1989, P 217-222 22. A.F. VOTER AND S.P. CHEN, ACCURATE INTERATOMIC POTENTIALS FOR NI, AL, AND NI3AL, MATER. RES. SOC. PROC., VOL 82, 1987, P 175-180 23. S.P. CHEN, ANOMALOUS RELAXATIONS OF (0001) AND (10 1 0) SURFACES IN HCP METALS, SURF. SCI. LETT., VOL 264, 1992, P L162-L168 24. R.J. HARRISON, A.F. VOTER, AND S.P. CHEN, ATOMISTIC MODELING OF MATERIALS: BEYOND PAIR POTENTIALS, V. VITEK AND D.J. SROLOVITZ, ED., PLENUM PRESS, 1989, P 219-222 25. R.J. HARRISON, F. SPAEPEN, A.F. VOTER, AND S P. CHEN, STRUCTURE OF GRAIN BOUNDARIES IN IRON, PROC. 34TH SAGAMORE ARMY MATERIALS RESEARCH CONFERENCE (LAKE GEORGE, NY), 1990, P 651-675 (FOR FE) 26. S.P. CHEN, LOCAL VOLUME POTENTIALS FOR BCC METALS, TO BE PUBLISHED (FOR FE, NB, CR, W, AND MO) 27. S.P. CHEN, A.M. BORING, R.C. ALBER, AND P.J. HAY, THEORETICAL STUDIES OF NI3AL AND NIAL WITH IMPURITIES, HIGH TEMPERATURE ORDERED INTERMETALLIC ALLOYS III, C.C. KOCH, C.T. LIU, N.S. STOLOFF, AND A.I. TAUB, ED., PROC. MATER. RES. SOC., VOL 133, 1989, P 149-154 28. F.R. DEBOER, R. BOOM, W.C.M. MATTENS, A.R. MIEDEMA, AND A.K. NIESSEN, COHESION IN METALS: TRANSITION METAL ALLOYS, NORTH-HOLLAND, 1989 Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Properties Affecting Adhesion Large impurities, such as bismuth, sulfur, phosphorus, and so on, tend to lower grain boundary adhesion, whereas small interstitial atoms, such as boron, carbon, and beryllium tend to increase cohesion in nickel, Ni3Al, and Ni3Si (Ref 1, 20, 21). Deviations from stoichiometry can either increase or decrease cohesion, depending on whether amounts of strong bonding elements, like nickel in Ni3Al, are increased or decreased. Although stoichiometry may slightly affect the ideal cohesion energy, its effect on the consumption of plastic work is potentially higher, as demonstrated in Co3Ti. Researchers (Ref 29) found that excess cobalt suppressed intergranular brittleness and rendered the material ductile, without adding strengtheners such as boron or beryllium. Sometimes impurity effects only work in off-stoichiometric compounds, such as boron in Ni3Al. Quantitative analysis of these kinds of effects have been attempted for impuritycontrolled fracture properties in iron, which has been studied extensively by Hondros and Seah (Ref 30). A simple linear formula linking cohesion and level of impurities has been extracted from many experiments (Ref 30).


1. S.P. CHEN, D.J. SROLOVITZ, AND A.F. VOTER, COMPUTER SIMULATION ON SURFACES AND [001] TILT GRAIN BOUNDARIES IN NI, AL, AND NI3AL, J. MATER. RES., VOL 4, 1989, P 62-77 20. S.P. CHEN, A.F. VOTER, R.C. ALBERS, A.M. BORING, AND P.J. HAY, INVESTIGATION OF BORON EFFECT ON NI AND NI3AL GRAIN BOUNDARIES, J. MATER. RES., VOL 5, 1990, P 955970 21. S.P. CHEN, A.F. VOTER, R.C. ALBERS, A.M. BORING, AND P.J. HAY, THEORETICAL STUDIES OF GRAIN BOUNDARIES IN NI3AL WITH BORON OR SULFUR, SCR. METALL., VOL 23, 1989, P 217-222 29. J.E. HACK, D.J. STROLOVITZ, AND S.P. CHEN, A MODEL FOR THE FRACTURE BEHAVIOR OF POLYCRYSTALLINE NI3AL, SCR. METALL., VOL 20, 1986, P 1699-1704

30. E.D. HONDROS AND M.P. SEAH, INTERFACIAL AND SURFACE MICROCHEMISTRY, PHYSICAL METALLURGY, 3RD ED., R.W. CAHN AND P. HAASEN, ED., ELSEVIER SCIENCE PUBLISHER, 1983, P 855-931 Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Adhesion Measurement Sikorski (Ref 31) describes two modifications of the twist-compression bonding method for adhesion testing and concludes that the measured coefficient of friction supports the adhesion theory of friction. This correlation implies that a large coefficient of adhesion (ratio of the force necessary to break a bond to the normal loading force used to produce the bond) corresponds to a large coefficient of friction. Experimentally, he showed that the coefficient of adhesion was inversely proportional to the melting point, which meant that higher melting point materials had lower adhesion. This agress with the observation that higher melting point materials have higher hardness, which also corresponds to lower adhesion coefficients. Gerkema and Miedema (Ref 32) considered lubrication of iron substrates by thin films of lead and silver with copper, platinum, or molybdenum additions by calculating surface adhesion using heats of solution and surface energies. Their predictions agree well with sliding wear experiments. Pashley and Tabor (Ref 33) measured adhesion between tungsten and nickel using techniques similar to those used in atomic force microscopy (AFM). They were able to verify surface cleanliness and ensure that the adhesion measured was between clean tungsten and nickel by using Auger spectroscopy (AES). In a subsequent publication, Pashley and Pethia again measured the surface energy of adhesion between tungsten and nickel (Ref 34). This work was done between a tungsten tip and a nickel flat in high vacuum and showed that surface forces alone could cause plastic flow. The loading/unloading curves provide information on the elastic/plastic response of the material. This in turn determines the interpretation of the adhesion. Johnson and Keller (Ref 35) used a torsion balance with a 3 mm (0.12 in.) diameter ball against a flat plate to measure the force of adhesion between silver-silver, copper-nickel, and silver-nickel couples. They concluded that mutual solubility did not affect the tendency for adhesion. However, surface contaminants did have a significant effect on the adhesion (joint strength) between silver-silver couples when the vacuum varied. The difference in joint strength for a clean surface at 130 μPa (5 × 10-10 torr) and for the same surface after a few microns of argon was 50 GPa (5 × 108 dynes/cm2). Conrad and Rice (Ref 36) studied the cohesion of fcc metals (silver, aluminum, copper, and nickel) by fracturing the metals in high vacuum, rewelding the metals by compressing the fractured surfaces, and retesting the metals by fracturing. They found that the cohesive strength of the weld increased with compressive load when both were normalized to the initial fracture strength. Results for all materials fell on a single curve. They proposed that the essentially constant cohesion coefficient was due to the fact that the contact area produced by a given compressive load was inversely proportional to the fracture stress of the metal, whereas the cohesive strength was directly proportional to the fracture stress. Acoustic emission was used to measure and monitor cohesion. In one case, this technique was used to signal the onset of fiber debonding and to determine crack initiation locations in fiber-reinforced ceramic-matrix composites (Ref 37). This approach is a valuable and viable technique for monitoring debonding, but cannot be used to provide quantitative information on either bond strength or fundamental material parameters. It can be used to locate the origin of the debonding, but the error is on the order of the dimensions used for continuum calculations, which are much larger than the atomic dimensions. In another study (Ref 38), acoustic measurements were used to monitor fatigue cracks and interfaces by defining a spring stiffness (for continuous and discontinuous spring models) that represented the degree of contact across an interface. A measured acoustic reflection coefficient was found to be in good agreement with one predicted for the same geometry. In addition, a single stress-intensity factor was found to determine bond strength.

In a similar vein, Neid (Ref 39) developed a technique for measuring bond integrity by propagating a planar shock were through the bond region. Depending on the shock energy, defects developed in the bond. The change in density was used to detect defects created at the interface, because other nondestructive test (NDT) methods were not sensitive enough to detect the small defects created. The defects were subsequently verified using optical microscopy. Itou (Ref 40) used electrical resistance to monitor crack tip contact in order to verify crack tip shielding in aluminum, steel, and cork. The results were inconclusive, although he found that there was no oscillating stress singularity around the crack tip. From this discussion, one can see that no generally quick, accurate, and easy-to-use technique has been developed to measure adhesion.


31. M.E. SIKORSKI, THE ADHESION OF METALS AND FACTORS THAT INFLUENCE IT, WEAR, VOL 7, 1964, P 144-162 32. J. GERKEMA AND A.R. MIEDEMA, ADHESION BETWEEN SOLID METALS: OBSERVATIONS OF INTERFACIAL SEGREGATION EFFECTS IN METAL FILM LUBRICATION EXPERIMENTS, SURF. SCI., VOL 124, 1983, P 351-371 33. M.D. PASHLEY AND D. TABOR, ADHESION AND DEFORMATION PROPERTIES OF CLEAN AND CHARACTERIZED METAL MICRO-CONTACTS, VACUUM, VOL 10-12, 1981, P 619-623 34. M.D. PASHLEY AND J.B. PETHICA, THE ROLE OF SURFACE FORCES IN METAL-METAL CONTACTS, J. VAC. SCI. TECHNOL. A, VOL 3 (NO. 3), MAY/JUNE 1985, P 757-761 35. K.I. JOHNSON AND D.V. KELLER, JR., EFFECT OF CONTAMINATION ON THE ADHESION OF METALLIC COUPLES IN ULTRA-HIGH VACUUM, J. APPL. PHYS., VOL 38 (NO. 4), JAN 1967, P 1896-1904 36. H. CONRAD AND L. RICE, THE COHESION OF PREVIOUSLY FRACTURED FCC METALS IN ULTRA-HIGH VACUUM, METALL. TRANS., VOL 1, NOV 1970, P 3019-3029 37. T. KISHI, M. ENOKI, AND H. TSUDA, INTERFACE AND STRENGTH IN CERAMIC MATRIX COMPOSITES, MATER. SCI. ENG., VOL A143, 1991, P 103-110 38. O. BUCK, R.B. THOMPSON, D.K. REHBEIN, D.D. PALMER, AND L.J.H. BRASCHE, CONTACTING SURFACES: A PROBLEM IN FATIGUE AND DIFFUSION BONDING, METALL. TRANS. A, VOL 20, APRIL 1989, P 627-635 39. H.A. NIED, "DETERMINATION OF THE STRUCTURAL INTEGRITY OF METALLURGICALLY BONDED PLATES BY HIGH ENERGY IMPULSE TEST," SIXTH INT. CONF. ON EXP. STRESS ANALYSIS (MUNICH), SEPT 1978 40. S. ITOU, AN EXPERIMENT ON THE BEHAVIOUR AROUND THE TIPS OF AN INTERFACE CRACK, ENG. FRACT. MECH., VOL 37 (NO. 1), 1990, P 145-150 Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Interfacial Analysis Electron spectroscopies are extremely valuable for the quantitative and qualitative analysis of surfaces. Fundamental information about the atomic forces and energies of the regions being probed, as well as their geometry, can be obtained. The shortcoming of these techniques in analyzing interfaces is that the interfacial volume is a small fraction of the probed volume and the effects of the interface cannot be separated from the signal. This condition can be overcome by using scanning tunneling microscopy (STM) or AFM. However, preparation of the surface is extremely difficult, the equipment requires a skilled and trained operator, the data are fundamental, and engineering information is not obtained. The crux of the problem is that there is no simple connection between condensed matter theory, micromechanics, and continuum

mechanics. The following discussion briefly reviews some of the available technical literature that describes the use of these spectroscopies and their application to bimaterial systems. Photoelectron spectroscopy has been used to study electron core-level shifts for ytterbium, aluminum, and silicon grown on molybdenum (110) (Ref 41, 42). The core-level shifts between the film and the substrate are expressed in terms of adhesion energy. In this case, the adhesion energy is the difference between the energy of the neutral and excited atoms per atom, which is readily found from the spectra. The measured adhesion energy is in good agreement with the adhesion energy determined using the Miedema technique (Ref 28). High-resolution electron microscopy (HREM), electron energy loss spectroscopy (EELS), microspectroscopy, and microdiffraction have been used to analyzed the chemical distributions on or near interface planes (Ref 43). This information can be used to support models requiring detailed stoichiometry, but does not yield adhesion values. Gerberich et al. (Ref 44) used selected-area-electron channeling patterns (SACP) and stereoimaging to quantify the strain in grains. The work was carried out under plane strain and plane stress conditions. The researchers determined strain distributions within 1 μm (40 μin.) of a stationary or growing crack in Fe-3Si single crystals with a {100} crack. Although this arrangement is not a bimaterial interface, the results reflect the effort needed to obtain fundamental materials information. Optical techniques, such as reflection (Ref 45), have also been used to probe interfaces and to aid in determining interfacial electronic structure. The data, in turn, are related to adhesion and segregation. For example, Lanning et al. (Ref 46) recently correlated reflectance spectra with surface free energy. This was then used as a measure of surface energy and the work of decohesion was calculated. The most promising technology for understanding interfaces is probably STM and associated techniques. Advances in the science and technology of scanning tunneling microscopy/spectroscopy (STM/S), AFM, photon scanning tunneling microscopy (PSTM), scanning chemical potential microscopy (SCPM), and related spectroscopies/microscopies offer new insight into interfaces (Ref 47). Low-energy electron diffraction (LEED), LEED-spot profile analysis (SPA-LEED), scanning low-energy electron microscopy (SLEEM), and photoemission electron microscopy (PEEM) are also being applied to surface analysis and adhesion. These technologies, although currently applied to the vacuum-surface interface, are appropriate tools for the study of grain boundaries and dissimilar metal surfaces. They provide excellent methods for measuring the electronic behavior at interfaces and, thus, the fundamental interactions that occur at these interfaces. This information must then be correlated to engineering properties predicted by micromechanical and continuum mechanics methods, as well as experiments.


28. F.R. DEBOER, R. BOOM, W.C.M. MATTENS, A.R. MIEDEMA, AND A.K. NIESSEN, COHESION IN METALS: TRANSITION METAL ALLOYS, NORTH-HOLLAND, 1989 41. N. MARTENSSON, A. STENBORG, O. BJORNEHOLM, A. NILSSON, AND J.N. ANDERSON, QUANTITATIVE STUDIES OF METAL-METAL ADHESION AND INTERFACE SEGREGATION ENERGIES USING PHOTOELECTRON SPECTROSCOPY, PHYS. REV. LETT., VOL 60 (NO. 17), APRIL 1988, P 1731-1734 42. J.N. ANDERSON, O. BJORNEHOLM, A. STENBORG, A. NILSSON, C. WIGREN, AND N. MARTENSSON, MEASUREMENT OF METAL-METAL ADHESION AND INTERFACE SEGREGATION ENERGIES BY CORE-LEVEL PHOTOELECTRON SPECTROSCOPY. AL AND SI ON MO(110), J. PHYS., CONDENS. MATTER, VOL 1, 1989, P 7309-7313 43. R.W. CARPENTER, HIGH RESOLUTION INTERFACE ANALYSIS, MATER. SCI. ENG., VOL A107, 1989, P 207-216 44. W.W. GERBERICH, D.L. DAVIDSON, AND M. KACZOROWSKI, EXPERIMENTAL AND THEORETICAL STRAIN DISTRIBUTIONS FOR STATIONARY AND GROWING CRACKS, J. MECH. PHYS. SOLIDS, VOL 38 (NO. 1), 1990 45. A. JOHNER AND P. SCHAAF, CALCULATION OF THE REFLECTION COEFFICIENTS OF INTERFACES: A SCATTERING APPROACH, PHYS. REV. B, VOL 42 (NO. 9), SEPT 1990, P 55165526

46. B.R. LANNING, T. FURTAK, AND G.R. EDWARDS, A METHOD FOR PREDICTION OF METALCERAMIC INTERFACIAL BOND ENERGIES, SUBMITTED FOR PUBLICATION IN J. MATER. RES. 47. R.J. COLTON, C.R.K. MARRIAN, AND J.A. STROSCIO, ED., PROC. FIFTH INT. CONF. SCANNING TUNNELING MICROSCOPY/SPECTROSCOPY (SAM '90) AND THE FIRST INT. CONF. NANOMETER SCALE SCIENCE AND TECHNOLOGY (NANO I) (BALTIMORE, MD), JUNE 1990, AND J. VAC. SCI. TECH. B, SECOND SERIES, VOL 9 (NO. 2), PART II, MARCH/APRIL 1991 Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Welding (Bonding) An extensive review of diffusion bonding is given by Gerken and Owczarski (Ref 48). This treatise reviewed the status of theory at that time (1965) and cited numerous examples. Eleven years later, Williams and Crossland (Ref 49, 50) highlighted cold and diffusion welding. They pointed out that the major technical concerns for these processes had not significantly changed in the interim. As recently as 1986, there was concern over the knowledge base of bonding and adhesion at interfaces (Ref 51). This group felt that detailed mechanisms of adhesion were largely missing, but that firstprinciple theories were beginning to appear. A strong need for collaboration between theorists and experimentalists was indicated. Furthermore, much information was still unknown about interfaces, and the field was ripe for expansion. This situation still applies. Described below are examples of some of the proposed models for bonding. Several mechanisms have been proposed to explain the formation of solid-state welds (Ref 52, 53). Mohamed and Washburn (Ref 52) proposed a two-stage mechanism that depends on breaking the surface oxide. Their work is supported by the work of Bay. (Ref 54), who developed a theoretical model for oxide breakage and bond formation on aluminum. Kawakatsu and Kitayama (Ref 53) describe bonding as being controlled by three metallurgical factors, which have been identified as bonds between solid-solution, two-phase, or intermetallic materials. Their studies involved bond-strength measurements and microstructural examination as a function of bonding conditions for copper-nickel, copper-silver, and iron-aluminum couples. They concluded that a minimum bonding temperature occurs at about the softening temperature of one of the materials. They also found that bond strength depends on recrystallization temperature. In the case of solidsolution materials, the interfaces tend to disappear after sufficient time above the recrystallization temperature, whereas this does not occur in two-phase and intermetallic materials. Excellent treatises on surface and interfacial segregation are provided by Dowben and Miller (Ref 55) and Hondros and Seah (Ref 30). The concept of migration or segregation of one component of a system to a surface or interface is considered in detail. This work provides insight into the behavior that must be considered when determining the possibility of adhesion between two dissimilar materials. Garmong et al. (Ref 56) describe the rate-limiting step in bonding as the complete elimination of porosity from the bond line and have developed a time-prediction model to support it. The model proceeds in two stages. First, long-wavelength asperities (roughness) are eliminated by plastic flow. Second, voids are eliminated by plastic flow and vacancy diffusion. This model is similar to that proposed by Hill and Wallach (Ref 57), who consider void elimination at the bond line. In this case, voids are created by surface asperities on the mating surfaces and pressure sintering is invoked for its elimination. The effect of grain size is also considered. Cahn (Ref 58) derived interfacial energy from thermodynamic considerations. He considered how the order parameter changes surface tension, elastic energy, and excess free energy when it affects the lattice parameter of an ordered material at an antiphase boundary. The work is not only appropriate for grain boundary interfaces in ordered materials, but is also relevant for dissimilar materials. Cline (Ref 59) reviews the difference between deformation and diffusion bonding. He defines diffusion bonding as not requiring significant deformation for bond formation. With respect to deformation bonding, Cline discusses the "film" and "energy barrier" theories and reviews the factors affecting the resultant bond. With respect to diffusion bonding, he defines two stages for bonding: microscopic deformation and diffusion. He then reviews the results for bonding aluminum and titanium and shows that both can be readily deformation bonded. The diffusion bonding of aluminum was simplified

by using a silver interlayer, whereas titanium required smooth, clean surfaces and protection from atmospheric contaminants. Derby and Wallach (Ref 60, 61) developed a thorough model for diffusion welding which accounts for: • • • • • • •

SURFACE DIFFUSION FROM SURFACE SOURCES TO A NECK VOLUME DIFFUSION FROM SURFACE SOURCES TO A NECK DIFFUSION ALONG THE BOND INTERFACE FROM INTERFACIAL SOURCES TO A NECK VOLUME DIFFUSION FROM INTERFACIAL SOURCES TO A NECK POWER-LAW CREEP DEFORMING THE RIDGE (SURFACE ROUGHNESS) PLASTIC YIELDING DEFORMING THE RIDGE VAPOR PHASE MASS TRANSPORT

Mass transport rate equations were developed in terms of the process variables (time, temperature, pressure, and material properties). The model is in reasonable agreement with experiment (Ref 62, 63). Guo and Ridley (Ref 64) consider the rate-limiting step in bonding to be the removal of interfacial voids that are due to the surface roughness and have developed a model to support this position. The main difference between their model and that of Derby and Wallach (Ref 60) is the void geometry used. Guo and Ridley use a lenticular-shaped cavity and have developed a model that agrees reasonably with the experimental data.


30. E.D. HONDROS AND M.P. SEAH, INTERFACIAL AND SURFACE MICROCHEMISTRY, PHYSICAL METALLURGY, 3RD ED., R.W. CAHN AND P. HAASEN, ED., ELSEVIER SCIENCE PUBLISHER, 1983, P 855-931 48. J.M. GERKEN AND W.A. OWCZARSKI, "A REVIEW OF DIFFUSION WELDING," WELDING RESEARCH COUNCIL BULLETIN NO. 109, OCT 1965 49. J.D. WILLIAMS AND B. CROSSLAND, PART 1 COLD PRESSURE WELDING, CME, MARCH 1976, P 65-67 50. J.D. WILLIAMS AND B. CROSSLAND, PART 2 DIFFUSION BONDING, CME, APRIL 1976, P 73-75 51. J.R. SMITH, PANEL REPORT ON INTERFACIAL BONDING AND ADHESION, MATER. SCI. ENG., VOL 83, 1986 52. H.A. MOHAMED AND J. WASHBURN, MECHANISM OF SOLID STATE PRESSURE WELDING, WELD. J. RES. SUPP., SEPT 1975, P 302S-310S 53. I. KAWAKATSU AND S. KITAYAMA, STUDY ON DIFFUSION BONDING OF METALS, TRANS. JIM, VOL 18, 1977, P 455-465 54. N. BAY, MECHANISMS PRODUCING METALLIC BONDS IN COLD WELDING, WELD. J. RES. SUPP., MAY 1983, P 137S-142S 55. P.A. DOWBEN AND A. MILLER, ED., SURFACE SEGREGATION PHENOMENA, CRC PRESS, 1990 56. G. GARMONG, N.E. PATON, AND A.S. ARGON, ATTAINMENT OF FULL INTERFACIAL CONTACT DURING DIFFUSION BONDING, METALL. TRANS. A, VOL 6, JUNE 1975, P 1269-1279 57. A. HILL AND E.R. WALLACH, MODELLING SOLID-STATE DIFFUSION BONDING, ACTA METALL., VOL 37 (NO. 9), 1989, P 2425-2437 58. J.W. CAHN, INTERFACIAL FREE ENERGY AND INTERFACIAL STRESS: THE CASE OF AN INTERNAL INTERFACE IN A SOLID, ACTA METALL., VOL 37 (NO. 3), 1989, P 773-776 59. C.L. CLINE, AN ANALYTICAL AND EXPERIMENTAL STUDY OF DIFFUSION BONDING, WELD. J., NOV 1966, P 481-S TO 489-S 60. D. DERBY AND E.R. WALLACH, THEORETICAL MODEL FOR DIFFUSION BONDING, MET. SCI.,

VOL 16, JAN 1982, P 49-56 61. D. DERBY AND E.R. WALLACH, DIFFUSION BONDING: DEVELOPMENT OF THEORETICAL MODEL, MET. SCI., VOL 18, SEPT 1984, P 427-431 62. D. DERBY AND E.R. WALLACH, DIFFUSION BONDS IN IRON AND A LOW-ALLOY STEEL, J. MATER. SCI., VOL 19, 1984, P 3149-3158 63. D. DERBY AND E.R. WALLACH, DIFFUSION BONDS IN COPPER, J. MATER. SCI., VOL 19, 1984, P 3140-3148 64. Z.X. GUO AND N. RIDLEY, MODELLING OF DIFFUSION BONDING OF METALS, MATER. SCI. TECH., VOL 3, NOV 1987, P 945-953 Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Interface Formation When considering the compression needed to bond dissimilar materials, the deformation-mechanism map work of Frost and Ashby (Ref 65) is helpful. The maps are in a stress/temperature space in which a single deformation mechanism is dominant. The stress used is a normalized tensile stress (tensile stress/shear modulus), and the temperature is the homologous temperature T/Tm, where T is the test temperature and Tm is the melting point of the material being tested, where both temperatures are expressed in degrees K. From these maps and the corresponding constitutive equations, contour plots of constant strain rate can be plotted. At this point, knowing any two values of strain rate, stress, or temperature, a point on the map will give the third variable and thee dominant deformation mechanism. This information can be a guide to determining the bonding mechanism when considering one of the previously discussed mechanisms that requires elimination of defects at the interface. Dammer et al. (Ref 66) describe a study where type 347 stainless steel was bonded to itself with several different interlayers. The interlayers were electrolytic-grade copper foil, electrolytic tough pitch (ETP) copper sheet, 80Au-20Cu foil, 72Ag-28Cu foil, and several combinations of silver, copper, and gold electroplates. They successfully bonded the specimens at temperatures below 540 °C (1005 °F), but found that the intermediate materials were essential to forming the weld. Bachin (Ref 67) discusses the interactions of metals and oxides when diffusion welding metals to metals and to ceramics. He discusses the reactions and exchanges at the interface by considering the formation of spinels and oxides. This leads to the conditions under which each is known to be operative and to the possibility of diffusion across the resulting interface. For example, Al + SiO2 proceeds by exchange:

4AL + 3SIO2 → 2AL2O3 + 3SI

(EQ 7)

When thin silicon dioxide is joined through a thin intermediate interlayer of aluminum, new phases precipitate at the dividing boundaries and the solid solution is decomposed when the limit of solubility is reached for the reaction products. This behavior results in high internal stresses and resultant microcracks. Caution is therefore required when welding these systems. A thorough review of the thermodynamics involved for the systems being considered is essential. Munir (Ref 68) developed equations that predict the amount of time required to dissolve oxides into their host metals for systems that have unstable oxides with respect to dissolution. He considered niobium, zirconium, titanium, copper, tantalum, cobalt, nickel, and iron and found reasonable agreement with the theory. Wang and Vehoff (Ref 69) discuss the segregation of atoms to interfaces. Their work deals with grain boundaries as interfaces, and they conclude that the cohesive energy of the interface at a fixed chemical potential is always less than the cohesive energy at a fixed concentration. The concept of segregation to the interface is probably important for bimaterial interfaces because of the different chemical potentials, diffusion rates, and, thus, concentrations that can exist. Because

these parameters are all temperature dependent, the bonding process and the service temperature will affect the interface segregation and, thus, its performance. This work provides a way to consider these effects.


65. H.J. FROST AND M.F. ASHBY, DEFORMATION MECHANISM MAPS, PERGAMON PRESS, 1982 66. P.A. DAMMER, R.E. MONROE, AND D.C. MARTIN, FURTHER STUDIES OF DIFFUSION BONDING BELOW 1000 °F, WELD. J., MARCH 1969, P 116-S TO 124-S 67. V.A. BACHIN, INTERACTION BETWEEN METALS AND THE OXIDES OF CERAMIC MATERIALS DURING DIFFUSION WELDING, AVT. SVARKA, NO. 10, 1982, P 15-17 68. Z.A. MUNIR, A THEORETICAL ANALYSIS OF THE STABILITY OF SURFACE OXIDES DURING DIFFUSION WELDING OF METALS, WELD. J., DEC 1983, P 333-S TO 336-S 69. J.-S. WANG AND H. VEHOFF, THE EFFECT OF THE MOBILITY OF SEGREGATED ATOMS ON INTERFACIAL EMBRITTLEMENT, SCR. METALL., VOL 25, 1991, P 1339-1344 Fundamentals of Metal and Metal-to-Ceramic Adhesion Ray Dixon and S.P. Chen, Los Alamos National Laboratory

Strength of Interfaces One of the most important and probably the most frequent questions: What is the strength of the interfacial bond? This question has often been answered by tensile testing. The difficulty is that the failure is often not at the interface and that it must be forced there. Additionally, there is concern that the strain is not one-dimensional across the bond and that the bond experiences triaxial strain in service. For these reasons, there has been considerable interests in determining fracture toughness and in detecting the failure of interfaces. Several reviews on the importance and extent of interfacial science were recently presented in one technical journal (Ref 70). Three reviews are noteworthy because they discuss the mechanical properties of interfaces (Ref 71, 72, 73). The efforts of Hutchinson and Suo (Ref 71) bring together much of the recent work on elastic fracture phenomena for bimetallic systems and builds on it to support work being done to understand mixed-mode failure in systems with interfaces. Shih (Ref 72) summarizes progress on interfacial mechanics and points out that fracture toughness, at various mode mixities, is needed to characterize fracture resistance of a weak plane in a composite. An application of this analysis has been demonstrated on the alumina/niobium metal-ceramic system by Stout et al. (Ref 74). Comninou (Ref 73) reviews the interface crack problem from the standpoint of open and closed crack tips and describes preliminary experimental results on fatigue and fracture of interface cracks in aluminum bonded by epoxy. Interface Strength Models. There has been considerable progress in analyzing the fracture behavior of interfaces with

various configurations. This work has led to the development of specimen geometries for systems with differing physical and mechanical properties. The analysis work is reviewed in this section and the experimental work is described in the section "Interface Strength Measurements" of this article. Suo and Hutchinson (Ref 75) have provided a complete solution to the semi-infinite crack lying along the interface of two infinite elastic layers under edge loads. This solution allows the determination of material properties for complex geometries, as well as the effect of welding two dissimilar materials. Jensen (Ref 76) demonstrates that shielding the full effect of mode II and III loading is important for interface cracks in brittle materials and that a proper interface fracture criteria needs to be modeled. He discusses a method for comparing different phenomenological fracture criteria and sets up micromechanical models describing shielding for interface geometries with kinks. The models of He and Hutchinson (Ref 77) explain conditions for interfacial crack propagation (kinking) into material on either side of the interface. Their models suggest that a crack propagates along an interface as long as the compliant material is tough and that the stiff material is at least as tough as the interface. They also show that the energy release rate is increased if the crack kinks into the more-compliant material and is diminished if it kinks into the stiff material.

Qu and Li (Ref 78) also analyzed cracks and dislocations at the interface between two anisotropic elastic half spaces. They obtained an explicit and real solution to the dislocation problem in the Stroh (Ref 79) formalism and showed that a planar interface between dissimilar anisotropic solids is completely characterized by no more than nine independent parameters. The work by Li and Ting (Ref 80, 81) is also relevant. Evans et al. (Ref 82, 83) concluded that fracture energy is not unique and is usually greater than the thermodynamic work of adhesion. The nonuniqueness is related to the mixity experienced by interface cracks. As a result, the fracture energy is influenced by the choice of test specimen. This result implies that when considering fracture energy for specific material combinations, the test geometry must be known, reported, and maintained. This also means that it is now possible to systematically study interfacial mechanical behavior and produce meaningful data on interface fracture toughness. Fischmeister et al. (Ref 84) modeled simple crack propagation, dislocation emission from crack tips, and twinning using atomistic calculations (molecular dynamics) embedded in a finite-element continuum. They also studied the behavior of cracks propagating in tungsten and iron on different planes and in different directions within these planes. Their results confirm that {100} should be the easiest cleavage plane in these body-centered cubic (bcc) tungsten and iron crystals. They further studied the propagation direction restrictions for cracks on particular planes. The modeling work was supported by experiments to verify the crack propagation behavior in WC-Co alloy, high-speed steel consolidated with nickel particles, and nickel coated on high-speed steel. Chen et al. (Ref 1, 9, 20, 21) studied crack growth in nickel, Ni3Al, Ni3Al plus boron, Ni3Al plus sulfur, and iridium grain boundaries using local volume potentials. They found that cohesive strength is important in determining grain boundary fracture and can be strongly affected by impurities and stoichiometry. Another approach to modeling fracture is given by Needleman (Ref 85), who used a cohesive zone-type interface model. His model takes full account of finite geometry changes and is used to study decohesion of a viscoplastic block from a rigid substrate. He found that the normal traction versus separation relation had an exponential form consistent with the universal binding energy relationship of Rose et al. (Ref 86) and of Ferrante and Smith (Ref 87). Ting (Ref 88) showed that oscillatory displacement at the interfacial crack surface in a bimaterial under a twodimensional deformation depends only on the material properties and is independent of specimen orientation. The exception is that bimaterials always produce oscillation and that there is an exception for one choice of relative orientation in one material combination. He also shows that for two different orthotropic materials, the oscillation depends on whether or not the Voigt elastic constants combinations, C11 + C22 + C12 , for the two materials are different. Cook (Ref 89) explains crack propagation threshold in brittle materials through the work done by the applied loading and that which is needed to create new surfaces during crack propagation. He showed that the mechanical energy release rate was a measure of the equilibrium surface energy of the material and the chemical environment. The latter affects the surface tension of both the fresh and aged crack. Reynolds et al. (Ref 90) investigated the difference between the concepts of cracks with zero normal and shear traction over the entire face of the crack and cracks with normal and shear traction specified at zero on the crack surface, except near the tip, where zero shear fraction and continuity of displacements normal to the crack were required. Of these two approaches, the first leads to oscillatory stress singularities and interpenetration of the material near the crack tip, whereas the second shows that a small region of contact exists at the tip, regardless, of the loading. The two methods require different treatments in terms of stress intensity, but predict the same strain-energy release rate. The work was done using the finite-element method (FEM) and, thus, also served to verify the formulation of each problem by FEM analysis. The results showed that the two formulations could be modeled by the absence or presence of interface elements that prevent interpenetration of the crack surfaces. Additionally, fracture parameters were calculated and compared well with previous analytic solutions. Interface Strength Measurements. Based on the theoretical analysis described above, as well as their own models,

the authors identified below have carried out excellent experimental work in measuring the mechanical properties of interfaces. Charalambides et al. (Ref 91) developed a test procedure to measure interface fracture resistance under mixed-mode conditions. Strain-energy release rates and stress-intensity factors can be obtained using this test. O'Dowd et al. (Ref 92) developed FEM model solutions and tested several geometries to determine interfacial fracture toughness as a function of mixity. They were able to explain near-tip contact and variation of mixity with distance. They also developed a procedure

for determining the effect of residual stress on the stress-intensity factor. Brady et al. (Ref 93) developed a buckled plate test to determine the transverse fracture toughness of unidirectional continuous-fiber composites. Their test was found to be independent of sample geometry and testing conditions. Mecholsky and Baker (Ref 94) developed a double cantilever beam specimen with a chevron notch to measure fracture toughness of ceramic-metal interfaces. The specimen is easy to make and introduces its own crack, and the test is easy to perform. Zhu and Byrne (Ref 95) looked at the effect of interfacial defects by bonding and subsequent tensile, fatigue, and photomicroscopy tests. They measured void ratios at the bond line and found the defects smaller than those detectable by current nondestructive test methods (~0.2 to 0.5 mm2 by 1 to 3 μm, or 3 × 10-4 to 8 × 10-4 in.2, or 40 to 120 μin., thick) could still significantly affect bond strength. They concluded that fatigue testing is a better test method than tensile testing because it is sensitive to a void ratio of approximately 27%, whereas tensile tests were not. Thouless (Ref 96) used epoxy-bonded glass slides to investigate interfacial failure under mixed-mode loading. He found that when the scale is sufficiently large, the sample could be considered homogeneous with a weak planar interface. This result produced a correlation between apparent fracture resistance and the degree of mixity. At a very detailed scale, the lack of homogeneity required that the crack-tip stress be determined by the position of the crack within the interface (epoxy). This condition leads to a change in failure mechanisms, whereby under pure mode-I conditions, the crack propagates entirely within the interfacial layer, whereas mode-II effects drive the crack toward the glass-epoxy interface. Nickel specimens with known surface roughnesses were diffusion welded under controlled conditions of temperature and pressure (Ref 97). The welds were interrupted at several different time intervals and the resulting interface was examined. The defect sizes were successfully related to surface tension and diffusion. It was also found that temperature increases produced a greater effect on reducing interfacial defects than did pressure. Meriin and Sliozberg (Ref 98) make the point that sound welds (bonds) cannot be obtained unless the voids are eliminated at the interface, which requires high temperatures and pressures. It further requires that mechanical tests be used as the principal detection method. This article has shown that there has been and continues to be active interest in interfacial behavior. The science is sufficiently complex that numerous approaches are useful to evaluate interfacial behavior, depending on the specific goal. However, it is obvious that the seemingly unrelated approaches all seek to understand the interface. It is therefore incumbent upon scientists in this field to be familiar with all developments concerning adhesion so that they can contribute to its understanding and future advancement.


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83. A.G. EVANS, M. RUHLE, B.J. DALGLEISH, AND P.G. CHARALAMBIDES, THE FRACTURE ENERGY OF BIMATERIAL INTERFACES, METALL. TRANS. A, VOL 21, SEPT 1990, P 2419-2429 84. H.F. FISCHMEISTER, H.E. EXNER, M.-H. POECH, S. KOHLHOFF, P. GUMBSCH, S. SCHMAUDER, L.S. SIGL, AND R. SPEIGLER, MODELING FRACTURE PROCESSES IN METALS AND COMPOSITE MATERIALS, Z. METALLKDE, BD 80, 1989, P 839-846 85. A. NEEDLEMAN, AN ANALYSIS OF TENSILE DECOHESION ALONG AN INTERFACE, J. MECH. PHYS. SOLIDS, VOL 38 (NO. 3), 1990, P 289-324 86. J.H. ROSE, J. FERRANTE, AND J.R. SMITH, UNIVERSAL BINDING ENERGY CURVES FOR METALS AND BIMETALLIC INTERFACES, PHYS. REV. LETT., VOL 47 (NO. 9), AUG 1981, P 675678 87. J. FERRANTE AND J.R. SMITH, THEORY OF THE BIMETALLIC INTERFACE, PHYS. REV. B, VOL 31 (NO. 6), MARCH 1985, P 3427-3434 88. T.C.T. TING, INTERFACE CRACKS IN ANISOTROPIC BIMATERIALS, J. MECH. PHYS. SOLIDS, VOL 38 (NO. 4), 1990, P 505-513 89. R.F. COOK, CRACK PROPAGATION THRESHOLDS: A MEASURE OF SURFACE ENERGY, J. MATER. RES., VOL 1 (NO. 6), NOV/DEC 1986, P 852-860 90. R.R. REYNOLDS, K. KOKINI, AND G. CHEN, THE MECHANICS OF THE INTERFACE CRACK USING THE FINITE ELEMENT METHOD, J. ENG. MATER. TECHNOL., VOL 112, JAN 1990, P 38-43 91. P.G. CHARALAMBIDES, J. LUND, A.G. EVANS, AND R.M. MCMEEKING, "A TEST FOR DETERMINING THE FRACTURE RESISTANCE OF BIMATERIAL INTERFACES," REPORT M87-4, MATERIALS DEPARTMENT, COLLEGE OF ENGINEERING, UC SANTA BARBARA 92. N.P. O'DOWD, C.F. SHIH, AND M.G. STOUT, TEST GEOMETRIES FOR MEASURING INTERFACIAL FRACTURE TOUGHNESS, INT. J. SOLIDS AND STRUCT., VOL 29, 1992, P 571-589 93. R.L. BRADY, R.S. PORTER, AND J.A. DONOVAN, A BUCKLED PLATE TEST TO MEASURE INTERFACIAL TOUGHNESS IN COMPOSITES, J. MATER. SCI., VOL 24, 1989, P 4138-4143 94. J.J. MECHOLSKY AND L.M. BAKER, A CHEVRON-NOTCHED SPECIMEN FOR FRACTURE TOUGHNESS MEASUREMENTS OF CERAMIC-METAL INTERFACES, CHEVRON-NOTCHED SPECIMENS: TESTING AND STRESS ANALYSIS, STP 855, J.H. UNDERWOOD, S.W. FREIMAN, AND F.I. BARATTA, ED., ASTM, 1984, P 324-336 95. X. ZHU AND J.G. BYRNE, EFFECT OF INTERFACE DEFECTS ON THE MECHANICAL PROPERTIES OF AUTOGENOUS DIFFUSION BONDED JOINTS, J. MATER. ENG., VOL 13, 1991, P 207-211 96. M.D. THOULESS, FRACTURE OF A MODEL INTERFACE UNDER MIXED-MODE LOADING, ACTA METALL. MATER., VOL 38 (NO. 6), 1990, P 1135-1140 97. V.S. GOSTOMEL'SKII, E.S. KARAKOZOV, AND A.P. TERNOVSKII, THE IMPORTANCE OF DIFFUSION AND SURFACE TENSION IN THE FORMATION OF CONTACT DURING DIFFUSION WELDING, AVT. SVARKA, NO. 4, 1980, P 28-31 98. A.V. MERIIN AND S.K. SLIOZBERG, SOME CHARACTERISTICS OF JOINTS IN THE SOLIDPHASE WELDING OF METALS, SVAR. PROIZ., NO. 5, 1973, P 22-23 Fundamentals of Friction Welding J.W. Elmer and D.D. Kautz, Lawrence Livermore National Laboratory

Introduction FRICTION WELDING (FRW) is a solid-state welding process in which the heat for welding is produced by the relative motion of the two interfaces being joined. This method relies on the direct conversion of mechanical energy to thermal energy to form the weld, without the application of heat from any other source. Under normal conditions no melting

occurs at the interface. Figure 1 shows a typical friction weld, in which a nonrotating workpiece is held in contract with a rotating workpiece under constant or gradually increasing pressure until the interface reaches the welding temperature. The rotational speed, the axial pressure, and the welding time are the principal variables that are controlled in order to provide the necessary combination of heat and pressure to form the weld. These parameters are adjusted so that the interface is heated into the plastic temperature range where welding can take place. Once the interface is heated, axial pressure is used to bring the weld interfaces into intimate contact. During this last stage of the welding process, atomic diffusion occurs while the interfaces are in contact, allowing a metallurgical bond to form between the two materials.

FIG. 1 SCHEMATIC SHOWING FUNDAMENTAL STEPS IN THE FRICTION WELDING PROCESS. (A) ONE WORKPIECE IS ROTATED, AND THE OTHER WORKPIECE IS HELD STATIONARY. (B) BOTH WORKPIECES ARE BROUGHT TOGETHER, AND AXIAL FORCE IS APPLIED TO BEGIN THE UPSETTING PROCESS. (C) WORKPIECE ROTATION IS STOPPED, AND THE UPSETTING PROCESS IS COMPLETED

Friction welding involves heat generation through friction abrasion (Ref 1, 2, 3), heat dissipation, plastic deformation, and chemical interdiffusion. The interrelation among these factors during FRW to complications when trying to develop predictive models of the friction welding process. However, from a qualitative standpoint the process is well understood through empirical FRW studies that have been performed on a wide variety of materials. Five qualitative factors influence the quality of a friction weld (Ref 4): • • • • •

RELATIVE VELOCITY OF THE SURFACES APPLIED PRESSURE SURFACE TEMPERATURE BULK MATERIAL PROPERTIES SURFACE CONDITION AND PRESENCE OF SURFACE FILMS

The first three factors are related to FRW, while the last two are related to the properties of the materials being joined. During FRW, the relative velocity, the applied pressure, and the duration of the force are the three variables that are controlled. The effect of these variables on weld quality will be discussed for the two basic friction welding processes: direct-drive welding and inertia-drive welding. The surface temperature is the critical parameter for ensuring good welds and is dependent on the processing conditions and the materials being joined. Although the surface temperature is not measured or directly controlled the effects of insufficient or excessive temperature are generally apparent through visual examination of the finished weld. The bulk material properties and the condition of the surfaces being joined affect both the frictional forces and the forging characteristics of the materials being joined. These factors will be discussed for the friction welding of both similar-material and dissimilar-material combinations.

Commercial FRW applications employ a number of variations on the basic FRW concepts. These variations were developed to accommodate different part geometries and to produce different metallurgical effects (Ref 5) and will not be discussed in this article.

References

1. T.H. HAZLET, FUNDAMENTALS OF FRICTION WELDING, SOURCE BOOK ON INNOVATIVE WELDING PROCESSES, AMERICAN SOCIETY FOR METALS, 1981, P 11-36 2. F.P. BOWDEN AND D. TABOR, THE FRICTION AND LUBRICATION OF SOLIDS, PART I, OXFORD UNIVERSITY PRESS, 1954 3. J. GODDARD AND H. WILMAN, A THEORY OF FRICTION AND WEAR DURING THE ABRASION OF METALS, WEAR, VOL 5, 1962 4. V.I. VILL, FRICTION WELDING OF METALS, TRANSLATED FROM THE RUSSIAN, AMERICAN WELDING SOCIETY, 1962 5. WELDING HANDBOOK, 8TH ED., VOL 2, AMERICAN WELDING SOCIETY, 1991, P 739-763 Fundamentals of Friction Welding J.W. Elmer and D.D. Kautz, Lawrence Livermore National Laboratory

Friction Welding Technology There are two principal FRW methods: direct-drive welding and inertia-drive welding. Direct-drive FRW, sometimes called conventional friction welding, uses a motor running at constant speed to input energy to the weld. Inertia-drive friction welding, sometimes called flywheel friction welding, uses the energy stored in a flywheel to input energy to the weld. These two FRW technologies produce inherently different metallurgical effects at the joint interface. Both FRW technologies can be applied through different types of relative motion in order to generate the friction necessary to form the weld. The most common FRW geometry is that shown in Fig. 1, in which one cylindrical component is held stationary and the other is rotated. However, in other methods, both components are rotated in opposite directions, or two-stationary components are pushed against a rotating piece positioned between them. Additional forms of FRW, such as radial, orbital, and linear reciprocating motions, have been developed for special part geometries. These alternate methods are discussed elsewhere (Ref 5, 6). Direct Drive Welding In direct-drive FRW, a machine resembling a lathe is equipped with a brake and clutch, a means of applying and controlling axial pressure, and a weld-cycle timer and displacement controller. The operation of a direct-drive machine consists of a friction phase where heat is generated, a stopping phase where the rotation is terminated, and a forging phase where the pressure is applied to join the pieces. The relationships among the process variables are shown in Fig. 2, which plots the rotational speed and the axial pressure as a function of time for typical weld. The time required to stop the spindle is also an important variable because it affects the weld temperature and the timing of the forging force.

FIG. 2 PLOT OF SELECTED PARAMETERS VERSUS TIME RELATIVE TO THE THREE PHASES OF THE DIRECTDRIVE FRW PROCESS

The forging phase starts at the instant when higher pressure (that is, a larger forging force) is applied in the weld cycle. Thus, the forging phase actually starts somewhere in the stopping phase. In general, the larger forging force can be applied (case 1) while the spindle is decelerating in the stopping phase (Fig. 2) or (case 2) after the spindle has stopped rotating at the end of the stopping phase. The difference in the two applications of the larger forging force is the presence of a second friction peak. In case 1, the torque will rise again to reach a second peak before dropping. This produces a torsional force. In case 2, especially when the stopping phase is very short due to rapid braking, frictional torque does not rise but actually starts to decrease at the onset of the forging phase. In this case, there is no torsional force, and forging is affected only by the upsetting force. The speed of rotation is the least sensitive process variable in that it can be varied over a wide range if heating time and pressure are properly controlled. For steels, the recommended peripheral velocity varies from 75 to 215 m/min (250 to 700 ft/min) (Ref 7). In general, higher speeds correspond to low weld heat inputs and are used to weld heat-sensitive materials such as hardenable steels. The friction force is generally applied gradually to the weld to help overcome the initial contact-torque peak. For carbon steels, a friction pressure of about 70 MPa (10 ksi) at the interface area is required to form a good joint. After the drive motor is disengaged from the workpiece, the forging force is applied to complete the weld. Typical forging forces for carbon steel are of 140 MPa (20 ksi) at the weld interface. Inertia-Drive Welding The inertia-drive FRW method uses a similar type of machine except that the spindle holding the rotating piece is attached to a flywheel. The flywheel controls the energy input to the weld. The moment of inertia of the flywheel is an important variable that is adjusted by adding or removing flywheels. The amount of energy stored in the flywheel is controlled by its speed. Once the spindle is at the correct speed, the drive system is disengaged, leaving a rotating flywheel mass. Axial pressure is then applied and held constant throughout the welding process. The applied pressure results in a decrease in the rotational speed, typically referred to as deceleration. In some cases, when the spindle has either nearly stopped or come to a complete stop, a higher forging force may be used. Figure 3 illustrates the inertia-drive FRW process, which is similar to the direct-drive method in that the weld typically takes place in two stages: friction and forging. However, some weld schedules do not require a forging stage.

FIG. 3 PLOT OF SELECTED PARAMETERS VERSUS TIME RELATIVE TO THE TWO PHASES OF THE INERTIADRIVE FRW PROCESS

The major difference between the direct-drive and inertia-drive methods is the friction speed. In inertia welding, the friction speed continuously decreases during the friction stage, while in direct welding the friction speed stays constant. The heat generated by the plastic deformation of materials at the faying surfaces, not the heat generated by friction in the friction phase, is the primary source in the forging phase of preventing rapid decrease of the temperature at the interface. The process variables that control the characteristics of an inertia weld are the flywheel size (moment of inertia), the flywheel speed, and the axial pressure. The weld energy is related to the first two variables and is a fixed quantity once they have been determined. The kinetic energy in the flywheel at any time during the weld is given by: E=

S ².1 C

where E is the energy (ft · lbf, or J), I is the moment of inertia (lb/ft2, or kg/m2), S is the flywheel speed (rev/min), and C is a conversion constant that is equal to 5873 for English units or 182.4 for metric units. The constant C is derived from: E=

1 mv ² 2

where v = ωr. Because ω= 2πs, v = 2πsr: 1 m(2π )² s ² r ² 2 1 E = mr ²(2π )² s ² 2 E=

Making mr2 = I, the previous equation becomes: E=

1 l (2π )² s ² 2

In SI units, the previous equation becomes:  m².Kg  1  1   min  E= = (2π )² s ²  l [ m².Kg ]      s²  2  min ²   60 s  1 (2π )² E= s ²l 2 60² Is ² E= C

2

Solving for C in the previous equation: C=

2*60² 60² = = 182.378 (2π ) 2 2π 2

The energy stored in the flywheel is proportional to its speed of rotation squared, S2. Therefore, a wide range of energy levels can be obtained without changing the flywheel to accommodate changes in part geometries. For large changes in the parts being joined, the capacity of an inertia welding machine can be modified by changing the flywheel moment of inertia if necessary. Flywheel Energy. The moment of inertia of the flywheel is selected to produce both the desired amount of kinetic

energy and the desired amount of forging. Forging results from the characteristic increase in torque that occurs at the weld interface as the flywheel slows and comes to rest. This increased torque, in combination with the axial pressure, produces forging. Because forging begins at some critical velocity, the amount of forging depends on the amount of energy remaining in the flywheel, which is a linear function of the flywheel moment of inertia. Large, low-speed flywheels produce greater forging force than small, high-speed flywheels even though they contain the same amount of kinetic energy. Although small, medium, and large amounts of flywheel energy produce similar heating patterns, the amount of energy greatly affects the size and shape of the weld upset, as shown in Fig. 4(a) for similar-metal joints.

FIG. 4 SCHEMATIC SHOWING EFFECT OF WELDING PARAMETERS ON THE FINISHED WELD NUGGET OBTAINED WHEN SIMILAR METALS ARE WELDED USING INERTIA-DRIVE FRW EQUIPMENT. (A) FLYWHEEL ENERGY. (B) INITIAL PERIPHERAL VELOCITY OF WORKPIECE. (C) AXIAL PRESSURE. SOURCE: REF 7

Peripheral Velocities. There is an optimum range of peripheral or linear surface velocities for each combination of

metals being joined. For welding steel to steel, the recommended initial peripheral velocity ranges from 90 to 460 m/min (300 to 1500 ft/min). However, welds can be made at velocities as low as 85 m/min (275 ft/min). Figure 4(b) shows the effect of initial peripheral velocity on weld shape for similar-metal welds. Axial Pressure Versus Peripheral Velocities. The effect of varying the axial pressure is opposite to the effect of

varying the velocity. Welds made at low axial pressure resemble welds made at medium velocity relative to the formation of weld upset and heat-affected zones (HAZ), as shown in Fig. 4(c). Excessive pressure produces a weld that has poor quality at the center and has a large amount of weld upset, similar to a weld made at a low velocity.


5. WELDING HANDBOOK, 8TH ED., VOL 2, AMERICAN WELDING SOCIETY, 1991, P 739-763 6. RECOMMENDED PRACTICES FOR FRICTION WELDING, ANSI/AWS C6. 1-89, AMERICAN WELDING SOCIETY, 1989

7. METALS HANDBOOK, 9TH ED., VOL 6, AMERICAN SOCIETY FOR METALS, 1983 Fundamentals of Friction Welding J.W. Elmer and D.D. Kautz, Lawrence Livermore National Laboratory

Metallurgical Parameters Friction welding can be used to join a wide range of similar and dissimilar materials. Metals, ceramics, metal-matrix composites (MMC), and polymers have all been joined by FRW, and many of the dissimilar-metal combinations that cannot be joined by conventional fusion welding techniques are readily joined by FRW methods. This section summarizes some of the metals that have been joined by FRW and discusses the metallurgical considerations that govern the properties of the resulting weld. Joining of Similar Metals The two general requirements for forming good friction welds are, first, that the materials to be joined can be forged and, second, that the materials can generate friction at the weld interface. The first requirement eliminates similar-metal welds in brittle materials such as ceramics, cast irons, and cemented carbides. However, ductile materials can sometimes be joined to these materials. The second requirements eliminates materials that contain alloying that provide dry lubrication. Free-machining additives to steel, graphite-containing alloys such as cast iron, and lead alloys may suffer from this requirement. Almost all other metal alloys can be welded to themselves by FRW techniques. Table 1 summarizes a number of common similar-metal weld joints that have been made using inertia-drive FRW. Metallic alloys known to form high-quality FRW joints include alloys based on aluminum, copper (copper-nickel, brass, bronze), iron (low-alloy steel, tool steel, stainless steel, maraging steel), nickel, titanium, tantalum, and many others (Ref 5, 6). Near full-strength metallurgical bonds can be produced for a very wide range of similar-metal alloy friction welds. The microstructure and mechanical properties of inertia-welded similar-metal joints for the following alloys can be found in the sources listed:

TABLE 1 PARAMETERS FOR INERTIA-DRIVE FRICTION WELDING OF TWO 25 MM (1 IN.) DIAM BARS MADE OF SIMILAR METALS WORK METAL

SPINDLE SPEED, REV/MIN

AXIAL FORCE KN LBF × 103

FLYWHEEL SIZE(A) KG · LB · M2 FT2

WELD ENERGY KJ FT · LBF × 103

METAL LOSS(B) MM IN.

TOTAL TIME(C), S

1018 STEEL 1045 STEEL 4140 STEEL INCONEL 718 MARAGING STEEL TYPE 410 STAINLESS TYPE 302 STAINLESS COPPER (COMMERCIALLY PURE) COPPER ALLOY 260 (CARTRIDGE BRASS, 70%) TITANIUM ALLOY, TI6AL-4V ALUMINUM ALLOY 1100

4600 4600 4600 1500 3000 3000 3500 8000

53 62 67 220 90 80 80 22

12 14 15 50 20 18 18 5

0.28 0.33 0.35 5.48 0.84 0.84 0.59 0.04

6.7 7.8 8.3 130.0 20.0 20.0 14.0 1.0

33 38 41 68 41 41 41 14

24 28 30 50 30 30 30 10

2.5 2.5 2.5 3.8 2.5 2.5 2.5 3.8

0.10 0.10 0.10 0.15 0.10 0.10 0.10 0.15

2.0 2.0 2.0 3.0 2.5 2.5 2.5 0.5

7000

22

5

0.05

1.2

14 10

3.8

0.15

0.7

6000

36

8

0.07

1.7

22 16

2.5

0.10

2.0

5700

27

6

0.11

2.7

20 15

3.8

0.15

1.0

ALUMINUM ALLOY 6061

5700

31

7

0.13

3.0

23 17

3.8

0.15

1.0

Source: Ref 7

(A) MOMENT OF INERTIA OF THE FLYWHEEL. (B) TOTAL SHORTENING OF THE WORKPIECES DURING WELDING. (C) SUM OF HEATING TIME PLUS WELDING TIME. The relative ease of friction welding metals to themselves is related to the matching properties at the weld interface. Because the materials properties are matched, heat is distributed uniformly on both sides of the joint, and the deformation characteristics are identical on both sides of the joint. This results in symmetric welds with good properties. In general, the process variables do not vary significantly for different alloys within a given class of materials. However, there can be a significant variation in processing variables between different classes of materials (Table 1). Because FRW generates localized heating at the interface, the HAZ is subject to rapid cooling due to heat transfer to the cold base metal. This rapid quenching may sufficiently alter the mechanical properties of the base metal in the HAZ region to require postweld heat treatment. For example, in order to restore ductility, stress relieving or tempering may be required to friction weld steels with hardenability greater than that of AISI 1035 (Ref 5, 6). In addition, age-hardenable alloys will lose strength in the HAZ during welding and may require postweld solution heat treating and/or postweld aging to restore their strength. Other alloys, such as those that obtain their strength from cold working, will lose strength in the HAZ of the weld, and their properties cannot be restored with postweld treatments. Joining of Dissimilar Metals While many similar-metal FRW joints are produced because of economic considerations, many dissimilar-metal FRW joints are produced because there are no alternative welding methods that can be used. Examples of these types of joints include dissimilar-metal combinations with widely different melting points and dissimilar-metal combinations that form incompatible phases when fusion welded. Table 2 gives parameters used for inertia welding several common dissimilarmaterial combinations.

TABLE 2 PARAMETERS FOR INERTIA-DRIVE FRICTION WELDING OF TWO 25-MM (1-IN.) DIAM BARS MADE OF DISSIMILAR METALS WORK METAL

SPINDLE SPEED, REV/MIN

AXIAL FORCE KN LBF × 103

FLYWHEEL SIZE(A) KG · LB · M2 FT2

WELD ENERGY KJ FT · LBF × 103

METAL LOSS(B) MM IN.

TOTAL TIME(C), S

Copper to 1018 steel M2 tool steel to 1045 steel Nickel alloy 718 to 1045 steel Type 302 stainless to 1020 steel Sintered high-carbon steel to 1018 Aluminum 6061 to type 302 stainless

8000 3000

22 180

5 40

0.06 1.14

1.4 27.0

20 54

15 40

3.8 2.5

0.15 0.10

1.0 3.0

1500

180

40

5.48

130.0

68

50

3.8

0.15

2.5

3000

80

18

0.84

20.0

41

30

2.5

0.10

2.5

4600

53

12

0.35

8.3

41

30

2.5

0.10

2.5

5500

22 5 (D) 67 15(D)

0.16 ...

3.9 ...

5.1 ...

0.20 ...

3.0 ...

Copper to aluminum alloy 1100

2000

33

0.46

11.0

27 20 . . ... . 10 7.5

5.1

0.20

1.0

7.5

(A) MOMENT OF INERTIA OF THE FLYWHEEL. (B) TOTAL SHORTENING OF THE WORKPIECES DURING WELDING.

(C) SUM OF HEATING TIME PLUS WELDING TIME. (D) THE 22 KN (5000 LBF) FORCE IS APPLIED DURING THE HEATING STAGE OF THE WELD; FORCE IS SUBSEQUENTLY INCREASED TO 67 KN (15000 LBF) NEAR THE END OF THE WELD. Low-Carbon Steels to Medium-Carbon Steels. In general, low- and medium-carbon steels are joined to each other

under a wide range of conditions, and high-carbon steels are readily joined to alloy steels (Ref 7) using friction welding. High-speed tool steels are welded to alloy steel shanks for numerous machine-tool applications. Steel with carbon contents as high as 1.0%, such as 52100 steel, can be joined to lower-carbon alloys. Preweld heat treating may be required in some cases to better match the properties at the interface, and postweld heat treatment may be required in some cases to temper the interface region of the high-carbon steel grades.

Stainless Steels to Other Selected Metals. Stainless steel alloys are comparatively easy to friction weld to other

metals. For example, austenitic stainless steel to low-alloy steel (Ref 15), titanium and copper to stainless steel (Ref 16), and 1100 aluminum to stainless steel (Ref 17) are examples of transition joints that are made by FRW. Titanium can be welded to stainless steel with extreme care (Ref 18), and other incompatible dissimilar combinations may be successfully welded using interlayer techniques (Ref 19). Figure 5 shows a micrograph of the interfacial region of an inertia weld between Monel 400 and 21-6-9 stainless steel. The joint properties are excellent, with plastic flow occurring in Monel 400 before joint failure during bend testing.

FIG. 5 METALLOGRAPHIC CROSS SECTION OF THE INTERFACE OF A MONEL 400 TO 21-6-9 STAINLESS STEEL WELD PRODUCED BY INERTIA-DRIVE FRW. NOTE THE FINE GRAIN SIZE PRESENT AT THE INTERFACE.

Such transition joints can often be used as interlayers for the friction welding of incompatible materials. For example, it is difficult to weld 5083 aluminum directly to stainless steel. However, by first friction welding aluminum alloy 1100 to the stainless steel, and machining the 1100 aluminum alloy back to an interlayer thickness of about 1 mm (0.04 in.), the 5083 aluminum alloy can be joined to the stainless steel via this 1100 interlayer with high joint efficiencies (Ref 20). Figure 6(a) shows an example of an aluminum-base MMC that was friction welded to 1100 aluminum. The MMC is a 2024 aluminum alloy with 15 vol% Al2O3 particles. The interface region between these materials is shown at higher magnification in Fig. 6(b), where intermixing of both the materials is shown to occur.

FIG. 6 CROSS SECTION OF THE INTERFACE OF A DIRECT-DRIVE FRICTION WELD JOINING 1100 ALUMINUM TO A 2024 ALUMINUM ALLOY WITH 15 VOL% AL2O3 PARTICLES. (B) HIGHER MAGNIFICATION OF THE SAME WELD SHOWING THE EXCELLENT WELD FORMED AT THE INTERFACE

Problems Common to Welding of Dissimilar Materials. In general, the same problems encountered when welding

similar materials must be addressed when welding dissimilar materials. However, some problems are associated only with the welding of dissimilar materials or are greatly magnified during the welding of dissimilar materials. These factors include joint interfaces, low-melting phases, brittle phases, and different thermal expansions. Joint Interfaces. While most similar-material welds are made with little concern for surface preparation, highly

dissimilar-metal combinations are more sensitive. This happens for various reasons. In stainless steel to aluminum alloy welds, the oxide surface that forms on the aluminum picks up contaminants such as water and hydrocarbons, forming extremely tenacious surface layers (Ref 21). If this layer is not removed prior to welding, poor structural welds may occur. In stainless steel to refractory metal alloy welds, the oxide on the faying surfaces again contain contaminants such as water and hydrocarbons. The contaminants in this case are likely to alloy into the finished weldment. This alloying causes a reduction of structural integrity through the formation of low-melting or brittle phases at the weld interface. Surface-treated interfaces frequently cause problems during FRW. Steels that have been carburized or nitrided, titanium alloys that have been nitrided, and other hardfaced alloys are difficult to friction weld due to the inherently low friction coefficient and low forgeability. The repeatability of welds made on materials with hard surface layers is difficult to characterize due to several factors, including coating thickness, coating quality, and physical properties of the coating. In most instances, weldability is improved if the surface-treated area is removed from the faying surface before welding. Low-Melting Phase Formation. Some material combinations have very low melting point phases associated with

mixing of constituents at the weld interface. The formation of these phases during the welding cycle is deleterious to the finished weld properties. Examples of combinations that fall into this category include iron-base alloys to titanium alloys and aluminum alloys to magnesium alloys. Low melting point eutectics are found in both of these metallurgical systems, and great care must be exercised during parameter development to prevent the formation of liquid phases during the completion of successful welds. Other weld combinations may be affected by contaminants at the weld interface. Examples include sulfur and phosphorus in iron-base alloys and bismuth in copper alloys. These contaminants may cause problems with hot shortness in very low concentrations. It is imperative that good cleaning practices be implemented when materials may have been contaminated with these elements or with material containing these elements. Brittle Phase Formation. Many materials, when combined, are susceptible to the formation of brittle phases. In some

combinations, this occurs during the welding cycle; in others, service conditions after welding cause the problem. Two main reasons exist for brittle phase formation in friction welds: • •

SURFACE CONTAMINANTS THAT EMBRITTLE THE WELD INTERFACE (SEE THE SECTION "JOINT INTERFACES" IN THIS ARTICLE) FORMATION OF INTERMETALLIC PHASES BETWEEN NORMAL CONSTITUENTS OF THE ALLOYS BEING WELDED

Intermetallic phase formation is common when welding refractory metal alloys to stainless steel alloys and in several other systems. In the case of stainless steels to refractory metals, σ phase or similar phases may occur upon welding at the interface. Proper weld procedure development reduces the amount of brittle phases that are formed, but typically does not eliminate their formation completely. Properly developed welds have satisfactory structural properties, because only small, noncontinuous areas of the brittle phase are present at the weld interface. Figure 7(a) shows an inertia-drive welded joint between vanadium and 21-6-9 stainless steel. The interface is smooth and shows no areas of brittle phases. Electron microscopy techniques are needed to find the small areas of σ phase present at the weld interfaces.

FIG. 7 METALLOGRAPHIC CROSS SECTION OF AN INERTIA-DRIVE FRW JOINT BETWEEN VANADIUM AND A 216-9 STAINLESS STEEL. NOTE THE EXCELLENT WELD QUALITY AT THE INTERFACE. (A) WELD INTERFACE WITH NO σ-PHASE GROWTH. (B) WELD INTERFACE WITH σ-PHASE GROWTH (INDICATED BY S) AND A SOLIDSOLUTION MIXING CAUSED BY CHEMICAL DIFFUSION AFTER EXPOSURE TO A 1000 °C (1830 °F) THERMAL CYCLE

Caution must be used when designing components for use at elevated-temperature extremes. In many instances, material combinations in which no brittle phases form during welding are susceptible to brittle phase formation at the interface during high-temperature use. This is not a design issue when the welds are used for near-room-temperature applications. Figure 7(b) shows the vanadium to 21-6-9 stainless steel inertia-drive weld after a severe thermal cycle of 1000 °C (1830 °F) for 2 h with σ-phase and solid-solution growth at the weld interface. The thick layer next to the stainless steel is a solid solution of iron and vanadium. The thin layer next to the vanadium is σphase and forms a continuous brittle fracture path across the weld interface. Differential Thermal Expansion. Some material combinations are difficult to weld because of the large differences in

thermal expansion. Low-expansion materials such as refractory metals, ceramics, and low-expansion iron-nickel and ironnickel-cobalt alloys may fail or be highly stressed during cooling when welded to high-expansion material such as austenitic stainless steels and nickel-base and cobalt-base superalloys. Use of these combinations requires the designer to consider the large stresses developed within the fabricated structure if the welds are restrained when exposed to large temperature changes. Intermediate expansion materials and multiple friction welds may be required to allow for the transition from high to low thermal expansion materials.


5. WELDING HANDBOOK, 8TH ED., VOL 2, AMERICAN WELDING SOCIETY, 1991, P 739-763 6. RECOMMENDED PRACTICES FOR FRICTION WELDING, ANSI/AWS C6. 1-89, AMERICAN WELDING SOCIETY, 1989 7. METALS HANDBOOK, 9TH ED., VOL 6, AMERICAN SOCIETY FOR METALS, 1983 8. T.N. HAZLET, PROPERTIES OF FRICTION WELDED PLAIN CARBON AND LOW ALLOY STEELS, WELD. J., VOL 41, 1962, P 49-S TO 52-S 9. S.B. DUNKERTON, TOUGHNESS PROPERTIES OF FRICTION WELDS IN STEELS, WELD. J., VOL 65, 1986, P 193-S TO 202-S 10. J.C. LIPPOLD AND B.C. ODEGARD, TECHNICAL NOTE: MICROSTRUCTURAL EVOLUTION

DURING INERTIA FRICTION WELDING OF AUSTENITIC STAINLESS STEELS, WELD. J., VOL 63, 1984, P 35-S TO 38-S 11. W.A. BAESLACK III AND K.S. HAGEY, INERTIA FRICTION WELDING OF RAPIDLY SOLIDIFIED POWDER METALLURGY ALUMINUM, WELD. J., VOL 67, 1988, P 139-S TO 149-S 12. H.H. KOO AND W.A. BAESLACK III, FRICTION WELDING OF A RAPIDLY SOLIDIFIED AL-FE-VSI ALLOY, WELD. J., VOL 71, 1992, P 147-S TO 169-S 13. K. OGAWA, H. YAMAGUCHI, T. MORIMOTO, K. TAKEMATA, H. SUDO, AND A. HIRATSUKA, SHEAR STRENGTH CHARACTERISTICS OF ALUMINUM ALLOY FRICTION WELDS, WELD. J., VOL 5, 1991, P 860-866 14. C.G. NESSLAR ET AL., FRICTION WELDING OF TITANIUM ALLOYS, WELD. J., VOL 50, 1971, P 379-S TO 385-S 15. K.G.K. MURTI AND S. SANDARESAN, THERMAL BEHAVIOR OF AUSTENITIC-FERRITIC JOINTS MADE BY FRICTION WELDING, WELD. J., VOL 64, 1985, P 327-S TO 334-S 16. R.A. BELL, J.C. LIPPOLD, AND D.R. ADOLPHSON, AN EVALUATION OF COPPER-STAINLESS STEEL FRICTION WELDS, WELD. J., VOL 63, 1984, P 325-2 TO 332-S 17. D. YASHAN, S. TSANG, W.L. JOHNS, AND M.W. DOUGHTY, INERTIA FRICTION WELDING OF 1100 ALUMINUM TO TYPE 316 STAINLESS STEEL, WELD. J., VOL 66, 1987, P 27-37 18. M. FUTAMATA AND A. FUJI, FRICTION WELDING OF TITANIUM AND SUS 304L AUSTENITIC STAINLESS STEEL, WELD. INT., VOL 4, 1990, P 768-774 19. F. SASSANI AND J.R. NEELAN, FRICTION WELDING OF INCOMPATIBLE MATERIALS, WELD. J., VOL 67, 1988, P 264-S TO 270-S 20. R. ARMSTRONG, PRIVATE COMMUNICATION AND INTERNAL REPORTS, LAWRENCE LIVERMORE NATIONAL LABORATORY, 1991 21. METALS HANDBOOK, 9TH ED., VOL 2, AMERICAN SOCIETY FOR METALS, 1979, P 204-209 Fundamentals of Friction Welding J.W. Elmer and D.D. Kautz, Lawrence Livermore National Laboratory

References

1. T.H. HAZLET, FUNDAMENTALS OF FRICTION WELDING, SOURCE BOOK ON INNOVATIVE WELDING PROCESSES, AMERICAN SOCIETY FOR METALS, 1981, P 11-36 2. F.P. BOWDEN AND D. TABOR, THE FRICTION AND LUBRICATION OF SOLIDS, PART I, OXFORD UNIVERSITY PRESS, 1954 3. J. GODDARD AND H. WILMAN, A THEORY OF FRICTION AND WEAR DURING THE ABRASION OF METALS, WEAR, VOL 5, 1962 4. V.I. VILL, FRICTION WELDING OF METALS, TRANSLATED FROM THE RUSSIAN, AMERICAN WELDING SOCIETY, 1962 5. WELDING HANDBOOK, 8TH ED., VOL 2, AMERICAN WELDING SOCIETY, 1991, P 739-763 6. RECOMMENDED PRACTICES FOR FRICTION WELDING, ANSI/AWS C6. 1-89, AMERICAN WELDING SOCIETY, 1989 7. METALS HANDBOOK, 9TH ED., VOL 6, AMERICAN SOCIETY FOR METALS, 1983 8. T.N. HAZLET, PROPERTIES OF FRICTION WELDED PLAIN CARBON AND LOW ALLOY STEELS, WELD. J., VOL 41, 1962, P 49-S TO 52-S 9. S.B. DUNKERTON, TOUGHNESS PROPERTIES OF FRICTION WELDS IN STEELS, WELD. J., VOL 65, 1986, P 193-S TO 202-S 10. J.C. LIPPOLD AND B.C. ODEGARD, TECHNICAL NOTE: MICROSTRUCTURAL EVOLUTION

DURING INERTIA FRICTION WELDING OF AUSTENITIC STAINLESS STEELS, WELD. J., VOL 63, 1984, P 35-S TO 38-S 11. W.A. BAESLACK III AND K.S. HAGEY, INERTIA FRICTION WELDING OF RAPIDLY SOLIDIFIED POWDER METALLURGY ALUMINUM, WELD. J., VOL 67, 1988, P 139-S TO 149-S 12. H.H. KOO AND W.A. BAESLACK III, FRICTION WELDING OF A RAPIDLY SOLIDIFIED AL-FE-VSI ALLOY, WELD. J., VOL 71, 1992, P 147-S TO 169-S 13. K. OGAWA, H. YAMAGUCHI, T. MORIMOTO, K. TAKEMATA, H. SUDO, AND A. HIRATSUKA, SHEAR STRENGTH CHARACTERISTICS OF ALUMINUM ALLOY FRICTION WELDS, WELD. J., VOL 5, 1991, P 860-866 14. C.G. NESSLAR ET AL., FRICTION WELDING OF TITANIUM ALLOYS, WELD. J., VOL 50, 1971, P 379-S TO 385-S 15. K.G.K. MURTI AND S. SANDARESAN, THERMAL BEHAVIOR OF AUSTENITIC-FERRITIC JOINTS MADE BY FRICTION WELDING, WELD. J., VOL 64, 1985, P 327-S TO 334-S 16. R.A. BELL, J.C. LIPPOLD, AND D.R. ADOLPHSON, AN EVALUATION OF COPPER-STAINLESS STEEL FRICTION WELDS, WELD. J., VOL 63, 1984, P 325-2 TO 332-S 17. D. YASHAN, S. TSANG, W.L. JOHNS, AND M.W. DOUGHTY, INERTIA FRICTION WELDING OF 1100 ALUMINUM TO TYPE 316 STAINLESS STEEL, WELD. J., VOL 66, 1987, P 27-37 18. M. FUTAMATA AND A. FUJI, FRICTION WELDING OF TITANIUM AND SUS 304L AUSTENITIC STAINLESS STEEL, WELD. INT., VOL 4, 1990, P 768-774 19. F. SASSANI AND J.R. NEELAN, FRICTION WELDING OF INCOMPATIBLE MATERIALS, WELD. J., VOL 67, 1988, P 264-S TO 270-S 20. R. ARMSTRONG, PRIVATE COMMUNICATION AND INTERNAL REPORTS, LAWRENCE LIVERMORE NATIONAL LABORATORY, 1991 21. METALS HANDBOOK, 9TH ED., VOL 2, AMERICAN SOCIETY FOR METALS, 1979, P 204-209 Fundamentals of Diffusion Bonding Murray W. Mahoney and Cliff C. Bampton, Rockwell International Science Center

Introduction DIFFUSION BONDING is only one of many solid-state joining processes wherein joining is accomplished without the need for a liquid interface (brazing) or the creation of a cast product via melting and resolidification (welding). In its most narrow definition, which is used to differentiate it from other joining processes such as deformation bonding or transient liquid phase joining, diffusion bonding (DB) is a process that produces solid-state coalescence between two materials under the following conditions: • • •

JOINING OCCURS AT A TEMPERATURE BELOW THE MELTING POINT, TM, OF THE MATERIALS TO BE JOINED (USUALLY >1/2TM ). COALESCENCE OF CONTACTING SURFACES IS PRODUCED WITH LOADS BELOW THOSE THAT WOULD CAUSE MACROSCOPIC DEFORMATION TO THE PART. A BONDING AID CAN BE USED, SUCH AS AN INTERFACE FOIL OR COATING, TO EITHER FACILITATE BONDING OR PREVENT THE CREATION OF BRITTLE PHASES BETWEEN DISSIMILAR MATERIALS, BUT THE MATERIAL SHOULD NOT PRODUCE A LOWTEMPERATURE LIQUID EUTECTIC UPON REACTION WITH THE MATERIALS TO BE JOINED.

Thus, diffusion bonding facilitates the joining of materials to produce components with no abrupt discontinuity in the microstructure and with a minimum of deformation.

Within the confines of this definition, the DB process, in practice, is limited to either press or gas pressure or bonding approaches. Illustrations and discussions of equipment and systems are presented elsewhere within this Handbook (see the Section "Solid-State Welding, Brazing, and Soldering Processes"). This article offers a qualitative summary of the theory of diffusion bonding. For those who require a more quantitative assessment, refer to the Selected References at the end of the article. It should be noted that the preferred term for this process, according to the American Welding Society, is diffusion welding. However, because diffusion bonding is used more commonly in industry, it is the term that will be used in this article. Fundamentals of Diffusion Bonding Murray W. Mahoney and Cliff C. Bampton, Rockwell International Science Center

Diffusion Bonding Process The DB process, that is, the application of pressure and temperature to an interface for a prescribed period of time, is generally considered complete when cavities fully close at the faying surfaces. Relative agreement is found for the mechanisms and sequence of events that lead to the collapse of interface voids, and the discussion below describes these metallurgical processes. Although this theoretical understanding of the DB process is universally applicable, it should be understood that parent metal strength is only approached for materials with surface conditions that do not have barriers to impede atomic bonding such as the absence of surface oxides or absorbed gases at the bonding interface. In practice, oxide-free conditions exist only for a limited number of materials. Accordingly, the properties of real surfaces limit and impede the extent of diffusion bonding. The most notable exception is titanium alloys, which, at DB temperatures greater than 850 °C (1560 °F), can readily dissolve minor amounts of adsorbed gases and thin surface oxide films and diffuse them away from the bonding surfaces, so that they will not impede the formation of the required metallic bonds across the bond interface, as shown in Fig. 1(a). An example of a successful diffusion bond in a titanium alloy is illustrated in Fig. 1(b), where the integrity of the bonded interfaces was demonstrated with subsequent superplastic expansion without interface failure.

FIG. 1 SUPERPLASTIC FORMING/DIFFUSION BONDING (SPF/DB) OF TITANIUM SHEET. (A) SEQUENCE OF OPERATIONS REQUIRED TO JOIN THREE SHEETS OF SUPERPLASTIC TITANIUM ALLOY USING SPF/DB PROCESS. (B) TYPICAL THREE-SHEET TITANIUM ALLOY COMPONENT SUPERPLASTICALLY FORMED FOLLOWING DIFFUSION BONDING.

Similarly, the joining of silver at 200 °C (390 °F) requires no deformation to break up and disperse oxides, because silver oxide dissociates completely at 190 °C (375 °F). Above this temperature, silver dissolves its oxide and also scavenges many surface contaminants. Other examples of metals that have a high solubility for interstitial contaminants include tantalum, tungsten, copper, iron, zirconium, and niobium. Accordingly, this class of alloy is easiest to diffusion bond. A second class of material, that is, metals and alloys that exhibit very low solubility for interstitials (such as aluminum-, iron-, nickel-, and cobalt-base alloys) are not readily diffusion bondable. Special consideration must be given to remove surface barriers to atomic diffusion prior to joining and subsequently prevent their reformation during the joining process. This is not an easy processing matter. Accordingly, the potential for high-strength bond interfaces for alloys with low interstitial solubility should be considered on an individual alloy basis. Fundamentals of Diffusion Bonding Murray W. Mahoney and Cliff C. Bampton, Rockwell International Science Center

Bonding Surfaces Containing Oxides Diffusion bonding can be achieved for materials with adherent surface oxides, but the resultant interface strengths of these materials are considerably less than that measured for the parent material. Aluminum alloys are prime examples of this class of material. Research since 1960 has demonstrated only limited diffusion bond properties. Although interface

strength can be increased for oxide-bearing materials, it requires considerable surface extension of the faying interfaces to create localized plastic flow of the metal and concurrent oxide breakage. This introduces an increased number of locations for metal-to-metal contact via plastic flow around or microextrusion through the broken oxide. In general, the oxide is not removed, but is simply dispersed over a greater surface area in an enclosed environment, in which oxidation cannot recur. Thus, even with significant surface deformation, only a fraction of the interface area contributes to the strength of the bond. The proportion of oxide-free metallic area revealed is dependent on the relative hardness of the metal and its oxide film, as well as on the mechanical properties of the oxide. This type of bonding, although often considered as diffusion bonding, is better described as deformation bonding and does not fit within the strict definition of the low deformation associated with diffusion bonding. Factors that affect the relative difficulty of diffusion bonding oxide-bearing surfaces include: • • •

•

SURFACE ROUGHNESS PRIOR TO WELDING. A ROUGHER SURFACE WILL RESULT IN GREATER SHEAR DEFORMATION. MECHANICAL PROPERTIES OF THE OXIDE. THE MORE BRITTLE THE OXIDE, THE GREATER THE DISPERSION FOR A GIVEN LEVEL OF DEFORMATION. RELATIVE HARDNESS OF THE METAL AND ITS OXIDE FILM. BECAUSE PLASTIC FLOW CONTROLS THE AMOUNT OF BONDING AREA, LARGE DIFFERENCES IN THEIR HARDNESS SHOULD FACILITATE BONDING. PRESTRAINING OR WORK HARDENING OF THE MATERIAL. INITIATION OF BONDING WILL OCCUR AT LOWER DEFORMATIONS FOR PRESTRAINED OR WORK-HARDENED MATERIALS, AND THE DEGREE OF SURFACE EXTENSION IN THE CENTRAL REGION OF THE INTERFACE IS CONSIDERABLY GREATER FOR COLD-WORKED MATERIAL. THUS, ANNEALED MATERIAL REQUIRES A LARGER TOTAL DEFORMATION BEFORE BONDING WILL INITIATE.

It is clear that with the appropriate information, sufficient experiments can be performed to determine the diffusion bondability of most materials. Parent metal strength will not always be attained using the DB approach, particularly for materials with adherent oxides, but interface strength can be maximized if the fundamentals of the process are understood. Fundamentals of Diffusion Bonding Murray W. Mahoney and Cliff C. Bampton, Rockwell International Science Center

Mechanism of Diffusion Bonding In diffusion bonding, the nature of the joining process is essentially the coalescence of two atomically clean solid surfaces. Complete coalescence comes about through a three-stage metallurgical sequence of events. Each stage, as shown in Fig. 2, is associated with a particular metallurgical mechanism that makes the dominant contribution to the bonding process. Consequently, the stages are not discretely defined, but begin and end gradually, because the metallurgical mechanisms overlap in time. During the first stage, the contact area grows to a large fraction of the joint area by localized deformation of the contacting surface asperities. Factors such as surface roughness, yield strength, work hardening, temperature, and pressure are of primary importance during this stage of bonding. At the completion of this stage, the interface boundary is no longer a planar interface, but consists of voids separated by areas of intimate contact. In these areas of contact, the joint becomes equivalent to a grain boundary between the grains on each surface. The first stage is usually of short duration for the common case of relatively high-pressure diffusion bonding.

FIG. 2 SEQUENCE OF METALLURGICAL STAGES IN DIFFUSION BONDING PROCESS. (A) INITIAL CONTACT: LIMITED TO A FEW ASPERITIES (ROOM TEMPERATURE). (B) FIRST STAGE: DEFORMATION OF SURFACE ASPERITIES BY PLASTIC FLOW AND CREEP. (C) SECOND STAGE: GRAIN BOUNDARY DIFFUSION OF ATOMS TO THE VOIDS AND GRAIN BOUNDARY MIGRATION. (D) THIRD STAGE: VOLUME DIFFUSION OF ATOMS TO THE VOIDS

During the second stage of joint formation, two changes occur simultaneously. All of the voids in the joints shrink, and most are eliminated. In addition, the interfacial grain boundary migrates out of the plane of the joint to a lower-energy equilibrium. Creep and diffusion mechanisms are important during the second stage of bonding and for most, if not all, practical applications, bonding would be considered essentially complete following this stage. As the boundary moves, any remaining voids are engulfed within grains where they are no longer in contact with a grain boundary. During this third stage of bonding, the voids are very small and very likely have no impact on interface strength. Again, diffusional processes cause the shrinkage and elimination of voids, but the only possible diffusion path is now through the volume of the grains themselves. Stage I: Microasperity Deformation

The nature of the starting surface is of considerable importance, because of the small macroscopic deformation allowed during diffusion bonding. A real surface is never perfectly clean or perfectly smooth, and the area of metal-to-metal contact between faying surfaces is a very small fraction of the area of joint contact. Contact is limited to a relatively few microasperities. At room temperature and under load, these asperities deform as long as the surface area of contact is such that the yield strength of the material is exceeded. The extent of this deformation is limited at room temperature, and is even more limited for work-hardenable materials. As temperature increases to the diffusion bonding temperature, the flow stress of the material decreases and additional asperity deformation occurs through plastic flow. Again, flow occurs until the area of contact increases to an extent that the yield strength of the material is exceeded. If the temperature is above the recrystalization temperature of the material, then work hardening is no longer a consideration. With time at temperature, creep mechanisms now control the rate of asperity deformation, and the area of contact or bond continues to grow. As the area of contract grows, the stress acting on the surface asperities decreases. Consequently, creep deformation progressively slows and diminishes in significance. The contributions of temperature and pressure to both plastic and creep deformation during this initial stage of diffusion bonding are synergistic, that is, at higher temperatures, less pressure is required and vice versa. However, for any combination of temperature and pressure, bulk deformation to the part is limited to a small percentage ( 5) than for the tensile specimens (εf ~0.001). The shorter rupture times for tensile joints are a direct result of the hydrostatic tensile stresses superimposed on the effective stresses caused by the mechanical constraint imposed by the base metal.

FIG. 9 COMPARISON BETWEEN THE RUPTURE TIME VERSUS EFFECTIVE (VON MISES) STRESS BEHAVIOR IN SHEAR, AT TORSIONAL STRESSES LESS THAN THE MAXIMUM OF HHC-DEPOSITED SILVER INTERLAYERS, COMPARED WITH THAT OF TENSION, CALCULATED USING FEM ANALYSES OF PM SPUTTER-DEPOSITED SILVER INTERLAYERS

The change in steady-state strain-rate (or time-to-rupture, according to the Monkman-Grant relationship) with applied stress is predictable according to the steady-state stress exponent, n,  ∂ ln ε  n=   ∂ ln σ 

(EQ 2)

where n is about 30 for silver at ambient temperature (Ref 38). Naturally, the steady-state strain rate of the torsional specimens shown in Fig. 9 is less than the constant applied strain rate of the torsional specimens shown in Fig. 8. If the steady-state rate if not achievable in the interlayer metal, then time-dependent failure in shear is not expected to occur. It is important to realize the consequence of the difference in plastic strains-to-failure between the constant shear-stress case and the constant tensile-stress case. For example, residual shear or tensile stresses may develop during fabrication of interlayer brazes or solid-state welds. This is particularly true in cases where dissimilar base materials are joined, but have substantially different coefficients of thermal expansion. For thin interlayers subjected to residual shear stresses, the dimensions of the base metal and interlayer may be such that the stress is relaxed at strains less than the shear strain-tofailure. However, because the observed tensile strains-to-failure are small, creep rupture by residual tensile stresses is possible, provided a significant stress level exists. If the interlayer metal is resistant to tim-dependent failure, then only the indirect implications of residual stress are important (for example, environmentally induced failure).


17. R.S. ROSEN, "TIME-DEPENDENT FAILURE OF SILVER INTERLAYER WELDS," PH.D.THESIS, UNIVERSITY OF CALIFORNIA, DAVIS, 1990 26. J.W. ELMER, M.E. KASSNER, AND R.S. ROSEN, THE BEHAVIOR OF SILVER-AIDED DIFFUSIONWELDED JOINTS UNDER TENSILE AND TORSIONAL LOADS, WELD. J., VOL 67 (NO. 7), 1988, P 157S-162S 33. M.E. KASSNER, R.S. ROSEN, AND G.A. HENSHALL, DELAYED MECHANICAL FAILURE OF SILVER-INTERLAYER DIFFUSION BONDS, METALL. TRANS. A, VOL 21, 1990, P 3085-3100 38. M.E. KASSNER, THE RATE DEPENDENCE AND MICROSTRUCTURE OF HIGH-PURITY SILVER DEFORMED TO LARGE STRAINS BETWEEN 0.16 AND 0.30 TM, METALL. TRANS. A, VOL 20, 1989, P 2001-2010 Mechanical Properties of Soft-Interlayer Solid-State Welds R.S. Rosen, Lawrence Livermore National Laboratory; M.E. Kassner, Oregon State University

Multiaxial Loading As discussed earlier, in tension there is significant mechanical constraint by the base material, which tends to reduce the effective stress. In torsion, where the shear stress is parallel to the plane of the interlayer, there is not constraint, other than, perhaps, the relatively minor constraint associated with plastic strain incompatibilities between the interlayer and the base metal. It was also mentioned earlier that the equivalent uniaxial strain-to-failure for torsional deformation of softinterlayer joints is roughly four orders of magnitude higher than the values for uniaxial tension, and the (effective) shear stress-to-failure is only about one-half the UTS. This suggests that the interlayers will behave anisotropically to imposed stresses. The mechanical properties of the intermediate cases of biaxial deformation are not readily interpolated between these limits. The authors have preliminary experimental evidence, for example, that low-strain ductile failures in thin interlayers may occur at particularly low stress levels and strains for cases where tensile loads are accompanied by inplane shear stresses (Ref 63). This implies that some multiaxial residual stress states may render solid-state-welded interlayer joints substantially more vulnerable to failure than simple tensile-stress states, even though the "macroscopic" effective stress may be comparable, when calculated on the basis of the applied stress state in the base material away from the interlayer.


63. M.E. KASSNER, OREGON STATE UNIVERSITY, UNPUBLISHED RESEARCH, 1989

Mechanical Properties of Soft-Interlayer Solid-State Welds R.S. Rosen, Lawrence Livermore National Laboratory; M.E. Kassner, Oregon State University

Environmentally Induced Failure of Interlayers The interlayer/base-metal interfaces of some soft-interlayer welds appear particularly vulnerable to failure when exposed to external (corrosive) environments. For example, it is known that brazed silver interlayers between stainless steel are subject to galvanic corrosion in aqueous NaCl media (Ref 64, 65). When dissimilar base materials are joined, the potential for stress-corrosion cracking (SCC) exists if the joints are exposed simultaneously to some external (corrosive) environments and either shear or tensile stresses. If an interlayer is subjected to stress, then a determination of the vulnerability to SCC is appropriate. It was mentioned earlier that elevated-temperature joining using either solid-state or brazing processes may leave the joint under substantial residual stress, particularly in the case of dissimilar metals for which the coefficients of thermal expansion are significantly different. The authors have reported (Ref 35) that the uranium to type 304 stainless steel solid-state welds, fabricated using HHC-deposited silver interlayers (Ref 25, 26), are very susceptible to SCC. These silver solid-state welds were joined at 873 K, and the difference in thermal expansion coefficients leaves the joints of cylindrical specimens under significant residual shear stresses after cooling to ambient temperature. It was found that these residual stresses led to SCC at the uranium/silver interface in air saturated with water vapor (100% relative humidity) (Ref 35). This cracking resulted in reduced tensile strengths and rupture times, when compared with those specimens tested in laboratory air ( 40% relative humidity). Specimens subjected to the same 100% relative humidity in air in an unstressed state exhibited no loss in strength, even after lengthy exposure. This, or course, emphasizes the importance of considering the residual stresses as not only rendering the interlayer joints vulnerable to "pure mechanical" rupture, but also to environmentally induced failures such as SCC, hydrogen embrittlement, and metal-induced embrittlement, if exposed to critical environments.


25. R.S. ROSEN, D.R. WALMSLEY, AND Z.A. MUNIR, THE PROPERTIES OF SILVER-AIDED DIFFUSION WELDS BETWEEN URANIUM AND STAINLESS STEEL, WELD. J., VOL 65 (NO. 4), 1986, P 83S-92S 26. J.W. ELMER, M.E. KASSNER, AND R.S. ROSEN, THE BEHAVIOR OF SILVER-AIDED DIFFUSIONWELDED JOINTS UNDER TENSILE AND TORSIONAL LOADS, WELD. J., VOL 67 (NO. 7), 1988, P 157S-162S 35. R.S. ROSEN, S. BEITSCHER, AND M.E. KASSNER, STRESS CORROSION CRACKING OF URANIUM-SILVER INTERFACES IN SILVER-AIDED DIFFUSION WELDS, INT. CONF. ENVIRONMENT-INDUCED CRACKING OF METALS, VOL 10, NACE, 1989, P 429-433 64. A.T. KUHN AND R.M. TRIMMER, REVIEW OF THE AQUEOUS CORROSION OF STAINLESS STEEL-SILVER BRAZED JOINTS, BR. CORROS. J., VOL 17 (NO. 1), 1982, P 4-8 65. T. TAKEMOTO AND I. OKAMOTO, EFFECT OF COMPOSITION ON THE CORROSION BEHAVIOR OF STAINLESS STEELS BRAZED WITH SILVER-BASE FILLER METALS, WELD. J., VOL 64, 1984, P 300S-307S Mechanical Properties of Soft-Interlayer Solid-State Welds R.S. Rosen, Lawrence Livermore National Laboratory; M.E. Kassner, Oregon State University

References

1. J.T. NIEMANN, R.P. SOPHER, AND P.J. RIEPPEL, DIFFUSION BONDING BELOW 1000 °F, WELD. J., VOL 37, 1958, P 337S-342S 2. I.M. BARTA, LOW TEMPERATURE DIFFUSION BONDING OF ALUMINUM ALLOYS, WELD. J., VOL 43, 1964, P 241S-247S

3. D. HAUSER, P.A. KAMMER, AND J.H. DEDRICK, SOLID-STATE WELDING OF ALUMINUM, WELD. J., VOL 46, 1967, P 11S-22S 4. P.A. KAMMER, R.E. MONROE, AND D.C. MARTIN, FURTHER STUDIES OF DIFFUSION BONDING BELOW 1000 °F, WELD. J., VOL 48, 1969, P 116S-124S 5. Y. IINO AND N. TAGAUCHI, INTERDIFFUSING METALS LAYER TECHNIQUE OF CERAMICMETAL BONDING, J. MATER. SCI. LETT., VOL 7 (NO. 9), 1988, P 981-982 6. A. URENA, J.M.G.D. SALAZAR, AND J. QUINONES, DIFFUSION BONDING OF ALUMINA TO STEEL USING SOFT COPPER INTERLAYER, J. MATER. SCI., VOL 27, 1992, P 599-606 7. T. YAMADA K, YOKOI, AND A. KOHNO, EFFECT OF RESIDUAL STRESS ON THE STRENGTH OF ALUMINA-STEEL JOINT WITH AL-SI INTERLAYER, J. MATER. SCI., VOL 25, 1990, P 21882192 8. M.G. NICHOLAS AND R.M. CRISPIN, DIFFUSION BONDING STAINLESS STEEL TO ALUMINA USING ALUMINUM INTERLAYERS, J. MATER. SCI., VOL 17, 1982, P 3347-3360 9. C.L. CLINE, AN ANALYTICAL AND EXPERIMENTAL STUDY OF DIFFUSION BONDING, WELD. J., VOL 45, 1966, P 481S-489S 10. A.T. D'ANNESSA, THE SOLID-STATE BONDING OF REFRACTORY METALS, WELD. J., VOL 43, 1964, P 232S-240S 11. M.G. NICHOLAS AND R.M. CRISPIN, DIFFUSION BONDING CERAMICS WITH DUCTILE METAL INTERLAYERS, PROC. INT. DIFFUSION BONDING (CRANFIELD, U.K.), 1987, P 173-182 12. B. DERBY, ZIRCONIA/METAL DIFFUSION BONDS, PROC. INT. CONF. DIFFUSION BONDING (CRANFIELD, U.K.), 1987, P 195-201 13. R.A. MORLEY AND J. CARUSO, THE DIFFUSION WELDING OF 390 ALUMINUM ALLOY HYDRAULIC VALVE BODIES, WELD. J., VOL 59 (NO. 8), 1980, P 29-34 14. C.H. CRANE, D.T. LOVELL, W.A. BAGINSKY, AND M.G. OLSEN, DIFFUSION WELDING OF DISSIMILAR METALS, WELD. J., VOL 46, 1967, P 23S-31S 15. M. O'BRIEN, C.R. RICE, AND D.L. OLSON, HIGH STRENGTH DIFFUSION WELDING OF SILVER COATED BASE METALS, WELD. J., VOL 55 (NO. 1), 1976, P 25-27 16. J.W. DINI, W.K. KELLEY, W.C. COWDEN, AND E.M. LOPEZ, USE OF ELECTRODEPOSITED SILVER AS AN AID IN DIFFUSION WELDING, WELD. J., VOL 63 (NO. 1), 1983, P 26S-34S 17. R.S. ROSEN, "TIME-DEPENDENT FAILURE OF SILVER INTERLAYER WELDS," PH.D.THESIS, UNIVERSITY OF CALIFORNIA, DAVIS, 1990 18. J.L. KNOWLES AND T.H. HAZLETT, HIGH-STRENGTH LOW-TEMPERATURE BONDING OF BERYLLIUM AND OTHER METALS, WELD. J., VOL 49 (NO. 7), 1970, P 301S-310S 19. P.S. MCLEOD AND G. MAH, THE EFFECT OF SUBSTRATE BIAS VOLTAGE ON THE BONDING OF EVAPORATED SILVER COATINGS, J. VAC. SCI. TECHNOL., VOL 11 (NO. 1), 1974, P 119-121 20. G. MAH, P.S. MCLEOD, AND D.G. WILLIAMS, CHARACTERIZATION OF SILVER COATINGS DEPOSITED FROM A HOLLOW CATHODE SOURCE, J. VAC. SCI. TECHNOL., VOL 11 (NO. 4), 1974, P 663-665 21. D.G. WILLIAMS, VACUUM COATING WITH A HOLLOW CATHODE SOURCE, J. VAC. SCI. TECHNOL., VOL 11 (NO. 1), 1974, P 374-377 22. E.R. NAIMON, D. VIGIL, J.P. VILLEGAS, AND L. WILLIAMS, ADHESION STUDY OF SILVER FILMS DEPOSITED FROM A HOY HOLLOW-CATHODE SOURCE, J. VAC. SCI TECHNOL., VOL 13 (NO. 6), 1976, P 1131-1135 23. E.R. NAIMON, J.H. DOYLE, C.R. RICE, D. VIGIL, AND D.R. WALMSLEY, DIFFUSION WELDING OF ALUMINUM TO STAINLESS STEEL, WELD. J., VOL 60 (NO. 1), 1981, P 17-20 24. D.T. LARSON AND H.L. DRAPER, CHARACTERIZATION OF THE BE-AG INTERFACIAL REGION OF SILVER FILMS DEPOSITED ONTO BERYLLIUM USING A HOT HOLLOW CATHODE DISCHARGE, THIN SOLID FILMS, VOL 107, 1983, P 327-334 25. R.S. ROSEN, D.R. WALMSLEY, AND Z.A. MUNIR, THE PROPERTIES OF SILVER-AIDED

DIFFUSION WELDS BETWEEN URANIUM AND STAINLESS STEEL, WELD. J., VOL 65 (NO. 4), 1986, P 83S-92S 26. J.W. ELMER, M.E. KASSNER, AND R.S. ROSEN, THE BEHAVIOR OF SILVER-AIDED DIFFUSIONWELDED JOINTS UNDER TENSILE AND TORSIONAL LOADS, WELD. J., VOL 67 (NO. 7), 1988, P 157S-162S 27. J. HARVEY, P.G. PATRIDGE, AND A.M. LURSHEY, FACTORS AFFECTING THE SHEAR STRENGTH OF SOLID STATE DIFFUSION BONDS BETWEEN SILVER-COATED CLAD AL-ZNMG ALLOY (ALUMINUM ALLOY 70100, MATER. SCI. ENG., VOL 79, 1986, P 191-199 28. P.G. PARTRIDGE AND D.V. DUNFORD, ON THE TESTING OF DIFFUSION-BONDED OVERLAP JOINTS BETWEEN CLAD AL-ZN-MG ALLOY (7010) SHEET, J. MATER. SCI. VOL 22, 1987 P 15971608 29. D.V. DUNFORD AND P.G. PARTRIDGE, THE PEEL STRENGTHS OF DIFFUSION BONDED. JOINTS BETWEEN CLAD AL-ALLOY SHEETS, J. MATER. SCI., VOL 22, 1987, P 1790-1798 30. R.S. ROSEN AND M.E. KASSNER, DIFFUSION WELDING OF SILVER INTERLAYERS COATED ONTO BASE METALS BY PLANAR-MAGNETRON SPUTTERING, J. VAC. SCI. TECHNOL. A, VOL 8 (NO. 1), 1900, P 19-29 31. M.E. KASSNER, R.S. ROSEN, G.A. HENSHALL, AND W.E. KING, DELAYED FAILURE OF SILVER AIDED DIFFUSION WELDS BETWEEN STEEL, PROC. 2ND INT. CONF. BRAZING, HIGH TEMPERATURE BRAZING, AND DIFFUSION WELDING, GERMAN WELDING SOCIETY, 1989, P 4752 32. M.E. KASSNER, R.S. ROSEN, G.A. HENSHALL, AND K.D. CHALLENGER, TIME-DEPENDENT FAILURE OF SILVER-INTERLAYER DIFFUSION WELDS BETWEEN ELASTICALLYDEFORMING BASE METALS, SCR. METALL. MATER., VOL 24, 1990, P 587-592 33. M.E. KASSNER, R.S. ROSEN, AND G.A. HENSHALL, DELAYED MECHANICAL FAILURE OF SILVER-INTERLAYER DIFFUSION BONDS, METALL. TRANS. A, VOL 21, 1990, P 3085-3100 34. P.G. PARTRIDGE AND J. HARVEY, "PROCESS FOR THE DIFFUSION BONDING OF ALUMINUM MATERIALS," U.K. PATENT APPLICATIONS GB 2117691 A, 1983 35. R.S. ROSEN, S. BEITSCHER, AND M.E. KASSNER, STRESS CORROSION CRACKING OF URANIUM-SILVER INTERFACES IN SILVER-AIDED DIFFUSION WELDS, INT. CONF. ENVIRONMENT-INDUCED CRACKING OF METALS, VOL 10, NACE, 1989, P 429-433 36. E. OROWAN, FRACTURE AND STRENGTH OF SOLIDS, REP. PROG. PHYS., VOL 12, 1948, P 185231 37. E. OROWAN, FUNDAMENTALS OF BRITTLE BEHAVIOR OF METALS, FATIGUE AND FRACTURE OF METALS, JOHN WILEY & SONS, 1952 38. M.E. KASSNER, THE RATE DEPENDENCE AND MICROSTRUCTURE OF HIGH-PURITY SILVER DEFORMED TO LARGE STRAINS BETWEEN 0.16 AND 0.30 TM, METALL. TRANS. A, VOL 20, 1989, P 2001-2010 39. N. BREDZ, INVESTIGATION OF FACTORS DETERMINING THE TENSILE STRENGTH OF BRAZED JOINTS, WELD. J., VOL 33, 1954, P 545S-563S 40. W.G. MOFFATT AND J. WULFF, TENSILE DEFORMATION AND FRACTURE OF BRAZED JOINTS, WELD. J., VOL 42, 1963, P 115S-125S 41. H.J. SAXTON, A.J. WEST, AND C.R. BARRETT, DEFORMATION AND FAILURE OF BRAZED JOINTS--MACROSCOPIC CONSIDERATIONS, METALL. TRANS., VOL 2, 1971, P 999-1007 42. A.J. WEST, H.J. SAXTON, A.S. TETELMAN, AND C.R. BARRETT, DEFORMATION AND FAILURE OF THIN BRAZED JOINTS--MICROSCOPIC CONSIDERATIONS, METALL. TRANS., VOL 2, 1971, P 1009-1017 43. E.P. LAUTENSCHLAGER, B.C. MARKER, B.K. MOORE, AND R. WILDES, STRENGTH MECHANISMS OF DENTAL SOLDER JOINTS, J. DENT. RES., VOL. 53, 1974, P 1361-1367 44. R.A. MUSIN, V.A. ANTSIVEROV, Y.A. BELIKOV, Y.V. LYAMIN, AND A.N. SOLOKOV, THE EFFECTS ON THE STRENGTH OF THE JOINT OF THE THICKNESS OF THE SOFT INTERLAYER

IN DIFFUSION WELDING, AUTO. WELD., VOL 32, 1979, P 38-40 45. R.M. TRIMMER AND A.T. KUHN, THE STRENGTH OF SILVER-BRAZED STEEL JOINTS--A REVIEW, BRAZING SOLDERING, VOL 2, 1982 P 6-13 46. N. BREDZ AND H. SCHWARTZBART, TRIAXIAL TENSION TESTING AND THE BRITTLE FRACTURE STRENGTH OF METALS, WELD. J., VOL 53 (NO. 12), 1956, P 610S-615S 47. W.G. MOFFATT AND J. WULFF, STRENGTH OF SILVER BRAZED JOINTS IN MILD STEEL, TRANS. AIME, VOL 209 (NO. 4), 1957, P 442-445 48. R.Z. SHRON AND O.A. BAKSHI, THE PROBLEMS OF GAUGING THE STRENGTHS OF WELDED JOINTS IN WHICH THERE IS A SOFT INTERLAYER, WELD. PROD., VOL 9, 1962, P 19-23 49. O.A. BAKSHI AND A.A. SHATOV, THE STRESSED STATE IN WELDED JOINTS WITH HARD AND SOFT INTERLAYERS, WELD PROD., VOL 13, 1966, P 13-19 50. V.S. GOLOVCHENKO AND B.A. GOLOBOV, THE USE OF SILVER AS AN INTERMEDIATE LAYER FOR JOINING TITANIUM TO OTHER METALS, WELD. PROD., VOL 18, 1971, P 55-57 51. G.K. KHARCHENKO AND A.I. IGNATENKO, THE STRENGTH OF JOINTS WITH A THIN SOFT INTERLAYER, AUTO. WELD., VOL 21, 1968, P 33-35 52. W.M. LEHER AND H. SCHWARTZBART, STATIC AND FATIGUE STRENGTHS OF METALS SUBJECTED TO TRIAXIAL STRESSES, VOL 60, 1960, P 610-626 53. E.A. ALMOND, D.K. BROWN, G.J. DAVIES, AND A.M. COTTENDED, THEORETICAL AND EXPERIMENTAL INTERLAYED BUTT JOINTS TESTED IN TENSION, INT. J. MECH. SCI., VOL 25 (NO. 3), 1983, P 175-189 54. G.A. HENSHALL, R.S. ROSEN, M.E. KASSNER, AND R.G. WHIRLEY, FINITE ELEMENT ANALYSIS OF INTERLAYER WELDS LOADED IN TENSION, WELD. J., VOL 69 (NO. 9) 1990, P 337S-345S 55. H.J. SAXTON, A.J. WEST, AND C.R. BARRETT, THE EFFECT OF COOLING RATE ON THE STRENGTH OF BRAZED JOINTS, METALL. TRANS., VOL 2, 1971, P 1019-1028 56. M.C. TOLLE AND M.E. KASSNER, TENSILE PROPERTIES OF THIN AU-NI BRAZES BETWEEN STRONG BASE METALS, SCR. METALL. MATER., VOL 26, 1992, P 1281-1284 57. M.E. KASSNER, M.C. TOLLE. R.S. ROSEN, G.A. HENSHALL, AND J.W. ELMER, RECENT ADVANCES IN UNDERSTANDING THE MECHANICAL BEHAVIOR OF CONSTRAINED THIN METALS IN BRAZES AND SOLID-STATE BONDS, THE METAL SCIENCE OF JOINING, M.J. CIESLAK, ED., TMS, 1992, P 223-232 58. R.W. LOGAN, R.G. CASTRO, AND A.K. MUKHERJEE, MECHANICAL PROPERTIES OF SILVER AT LOW TEMPERATURES, SCR. METALL., VOL 17, 1983, P 63-66 59. T.J. MOORE AND K.H. HOLKO, SOLID-STATE WELDING OF TD-NICKEL BAR, WELD. J., VOL 49 (NO. 9), 1970, P 395S-409S 60. R.S. ROSEN, "AN INVESTIGATION OF THE PROPERTIES OF A SILVER-AIDED SOLID-STATE BOND BETWEEN URANIUM AND STAINLESS STEEL," REPORT UCRL-53458, LAWRENCE LIVERMORE NATIONAL LABORATORY, 1983 61. E.R. NAIMON, R.G. KURZ, D. VIGIL, AND L. WILLIAMS, "SILVER FILMS FOR SOLID STATE BONDING,"REPORT RFP-3125, ROCKWELL INTERNATIONAL, 1981 62. Z. NISENHOLTZ, J. MIRONI, AND N. NIR, DIFFUSION BONDING OF STAINLESS STEEL 304L TO TI-6AL-4V ALLOY, PROC. 3RD INT. CONF. ISOSTATIC PRESSING (LONDON, U.K.), 1986 63. M.E. KASSNER, OREGON STATE UNIVERSITY, UNPUBLISHED RESEARCH, 1989 64. A.T. KUHN AND R.M. TRIMMER, REVIEW OF THE AQUEOUS CORROSION OF STAINLESS STEEL-SILVER BRAZED JOINTS, BR. CORROS. J., VOL 17 (NO. 1), 1982, P 4-8 65. T. TAKEMOTO AND I. OKAMOTO, EFFECT OF COMPOSITION ON THE CORROSION BEHAVIOR OF STAINLESS STEELS BRAZED WITH SILVER-BASE FILLER METALS, WELD. J., VOL 64, 1984, P 300S-307S

Shielded Metal Arc Welding Raymond H. Juers, Naval Surface Warfare Center

Introduction SHIELDED METAL ARC WELDING (SMAW), commonly called stick, or covered electrode, welding, is a manual welding process whereby an arc is generated between a flux-covered consumable electrode and the workpiece. The process uses the decomposition of the flux covering to generate a shielding gas and to provide fluxing elements to protect the molten weld-metal droplets and the weld pool. Shielded Metal Arc Welding Raymond H. Juers, Naval Surface Warfare Center

The SMAW Process The important features of the SMAW process are shown in Fig. 1. The arc is initiated by momentarily touching or "scratching" the electrode on the base metal. The resulting arc melts both the base metal and the tip of the welding electrode. The molten electrode metal/flux is transferred across the arc (by arc forces) to the base-metal pool, where it becomes the weld deposit covered by the protective, less-dense slag from the electrode covering.

FIG. 1 SMAW PROCESS

Advantages and Limitations. The SMAW process is the most widely used welding process. It is the simplest, in terms

of equipment requirements, but it is, perhaps, the most difficult in terms of welder training and skill-level requirements. Although welder skill level is a concern, most welders entering the field start as "stick welders" and develop the necessary skills through training and experience. The equipment investment is relatively small, and welding electrodes (except the very reactive metals, such as titanium, magnesium, and others) are available for virtually all manufacturing, construction, or maintenance applications. Shielded metal arc welding has the greatest flexibility of all the welding processes, because it can be used in all positions (flat, vertical, horizontal, and overhead), with virtually all base-metal thicknesses (1.6 mm, or

1 in., and greater), and in areas of limited accessibility, which is a very important capability. 16

Because the SMAW process is basically a manual process, the skill level of the welder is of paramount importance in obtaining an acceptable weld. The welder duty cycle is generally low, because of the built-in work break, which occurs after each electrode is consumed and requires replacement. In addition to replacing the electrode when the arc is stopped (broken), the welder may "chip" or remove slag and clean it away from the starting and welding area with a wire brush to allow the proper deposition of the subsequent weld. This electrode replacement and cleaning operation occurs many times during the work day (about every two minutes, or the time it generally takes to consume an electrode). This stopping,

chipping, wire brushing, and electrode replacement prevents the welder from attaining an operator factor, or duty cycle, that is much greater than 25%. Weld Quality. The quality of the weld depends on the design and accessibility of the joint, as well as on the electrode, the technique, and the skill of the welder. If joint details vary greatly from established design details, then a lower-quality weld can result. Other factors that also reduce quality are improper interbead cleaning, poor location of individual weld beads within the joint, and various problems with individual electrodes, including partially missing flux and core wires that are not centered within the flux covering. Overall, welds of excellent quality can be obtained with the SMAW process, as demonstrated by its use in joining submarine pressure hull sections and high-pressure oil/gas pipe lines. The base-metal thicknesses. that can be welded using the SMAW process generally range from 1.6 mm (

1 in.) to an 16

unlimited thickness. The thinner materials require a skilled welder, tight fitup, and the proper small-diameter welding electrode. Welding position also is important when determining the minimum plate thicknesses that can be welded. Flatposition butt welds and horizontal fillet welds are generally considered the easiest to weld. Out-of-position welding (vertical, overhead) requires greater skill. Welding Circuit. The circuit diagram for the SMAW process is shown in Fig. 2. The equipment consists of a power

source, electrode holder, and welding cables that connect the power source to the electrode holder and the workpiece. Alternating current (ac), or direct current, electrode negative (DCEN), or direct current, electrode positive (DCEP) can be used, depending on the electrode coating characteristics. The DCEN source is also called dc straight polarity, whereas the DCEP source is also called dc reverse polarity.

FIG. 2 SHIELDED METAL ARC CIRCUIT DIAGRAM

Equipment. The welding machine, or power source, is the crux of the SMAW process. Its primary purpose is to provide electrical power of the proper current and voltage to maintain a controllable and stable welding arc. Its output characteristics must be of the constant current (CC) type. SMAW electrodes operate within the range from 25 to 500 A. The electrode producer should suggest a narrow optimum range for each size and type of electrode. Operating arc voltage varies between 15 and 35 V.

The electrode holder, which is held by the welder, firmly grips the electrode and transmits the welding current to it. Electrode holders are available in several designs, such as the pincher type and the collet, or twist, type, shown in Fig. 3. Each style has its proponents and the selection is usually a personal preference. Electrode holders are designated by their current capacity. Selection factors, such as the current rating, duty cycle, maximum electrode size, and cable size, are shown in Table 1. The most lightweight holder that will accommodate the required electrode size is usually desired.

TABLE 1 SIZE AND CAPACITY OF ELECTRODE HOLDERS ELECTRODE HOLDER

RATING

MAXIMUM

MAXIMUM

ELECTRODE SIZE

CABLE SIZE

50

mm 3.2

1

200

50

4.0

MEDIUM

300

60

5.5

LARGE

400

60

6.4

EXTRA LARGE

500

75

7.9

600

75

9.5

SMALL

MAXIMUM CURRENT, A

DUTY CYCLE, %

100

in. 1 8 5 32 7 32 1 4 5 16 3 8

1/0 2/0 3/0 4/0 4/0

FIG. 3 SMAW ELECTRODE HOLDERS

All electrode holders should be fully insulated. Because they are used in proximity to the arc and are exposed to high heat, they will deteriorate rapidly. It is extremely important to maintain electrode holders to ensure that they retain their current-carrying efficiency, their insulating qualities, and their electrode gripping action. Manufacturers supply spare parts so that the holders can be rebuilt and maintained for safe and efficient operation. Certain pieces of auxiliary equipment can be used with the SMAW process, such as low-voltage control circuits, which enable the relatively high open-circuit voltage to be cut off until the electrode touches the workpiece. Other items include remote-control switches for the contactors, remote-control current adjusting devices, and engine idling controllers for engine-driven power sources. Shielded Metal Arc Welding Raymond H. Juers, Naval Surface Warfare Center

Applications Most manufacturing operations that require welding will strive to utilize the mechanized processes that offer greater productivity, higher quality, and, therefore, more cost-effective production. For these reasons, the SMAW process has been replaced where possible. However, the simplicity and ability of the SMAW process to achieve welds in areas of restricted accessibility means that it still finds considerable use in certain situations and applications. Heavy construction, such as shipbuilding, and welding "in the field," away from many support services that would provide shielding gas, cooling water, and other necessities, rely on the SMAW process to a great extent. Although the SMAW process finds wide application for welding virtually all steels and many of the nonferrous alloys, it is primarily used to join steels. This family of materials includes low-carbon or mild steels, low-alloy steels, high-strength steels, quenched and tempered steels, high-alloy steels, stainless steels, and many of the cast irons. The SMAW process is also used to join nickel and its alloys and, to a lesser degree, copper and its alloys. It can be, but rarely is, used for welding aluminum.

In addition to joining metals, the SMAW process is frequently used for the protective surfacing of base metals. The surfacing deposit can be applied for the purpose of corrosion control or wear resistance (hard surfacing). Shielded Metal Arc Welding Raymond H. Juers, Naval Surface Warfare Center

Electrodes The electrodes used in the SMAW process have many different compositions of core wire and a wide variety of fluxcovering types and weights. Standard electrode diameters of the core wire range from 1.6 to 8 mm (

1 5 to in.). 16 16

Electrode length usually ranges from 230 to 455 mm (9 to 18 in.); the shorter lengths are associated with the smallerdiameter electrodes. A bare, uncoated end of the electrode (the grip end) is provided for making electrical contact in the electrode holder. The coating on the electrode has numerous functions. It provides: •

• • • • •

GAS (NORMALLY, CARBON DIOXIDE), FROM THE DECOMPOSITION OF CERTAIN COATING INGREDIENTS TO SHIELD THE ARC AND WELD ZONE FROM THE ATMOSPHERE DEOXIDIZERS, FOR SCAVENGING AND PURIFYING THE DEPOSITED WELD METAL SLAG FORMERS, TO PROTECT THE DEPOSITED WELD METAL FROM ATMOSPHERIC OXIDATION AND TO HELP SHAPE THE WELD BEAD IONIZING ELEMENTS, TO MAKE THE ARC MORE STABLE AND TO OPERATE WITH ALTERNATING CURRENT ALLOYING ELEMENTS, TO PROVIDE SPECIAL CHARACTERISTICS TO THE WELD DEPOSIT IRON POWDER, IN CERTAIN ELECTRODES, TO INCREASE PRODUCTIVITY FOR WELDING FERROUS METALS

The American Welding Society (AWS) has established a system for identifying and classifying the different types of welding electrodes. All SMAW electrodes have the prefix letter E to indicate welding electrode. The symbols that follow the prefix are based on criteria that best describe the welding capabilities of the electrode metal. These criteria include chemical composition of the deposited weld metal, weld-metal mechanical properties, certain process parameters, or combinations of all factors. Mild and Low-Alloy Steel-Covered Electrodes. The prefix used to identify these electrodes is followed by a number series that indicates minimum strength level, position capability, and type of covering and welding current. Table 2 explains how the number series is used in AWS A5.1, the specification for carbon steel electrodes for shielded metal arc welding and AWS A5.5, the specification for low-alloy steel electrodes. The first two digits after the E in the E6010 electrode designate a tensile strength of at least 430 MPa (62 ksi) for the deposited metal in the as-welded condition. The third digit indicates the position in which satisfactory welds can be made with the electrode. Thus, the 1 in E6010, for example, means that the electrode is satisfactory for use in all positions (flat, vertical, horizontal, and overhead). The 2 in E6020 indicates that the electrode is suitable for the flat position and horizontal fillets. The last digit or last two digits, taken together, indicate the applicable current type to be used and the type of covering on the electrode.

TABLE 2 CARBON AND LOW-ALLOY STEEL COVERED ELECTRODE IDENTIFICATION SYSTEM

AWS MINIMUM TENSILE MINIMUM YIELD MINIMUM STRENGTH CLASSIFICATION(A) STRENGTH ELONGATION, % MPA KSI MPA KSI

E70XX E80XX E90XX E100XX E110XX E120XX CLASSIFICATION(B) EXXIX EXX2X EXX4X

480-500 550 620 690 760 830 FLAT POSITION YES YES YES

CLASSIFICATION(C)

CURRENT

ARC

EXX10 EXXX1

DCEP AC AND DCEP AC AND DCEN AC AND DC AC AND DC DCEP AC OR DCEP AC OR DCEP AC OR DC AC OR DC AC OR DC AC OR

EXXX2 EXXX3 EXXX4 EXXX5 EXXX6 EXXX8 EXX20 EXX24 EXX27 EXX28

70-72 390-420 80 460-550 90 530-620 100 600 110 670-760 120 740-830 HORIZONTAL POSITION YES FILLET YES PENETRATIONN

57-60 67-80 77-90 87 97-110 107-120 VERTICAL POSITION YES NO DOWN

17-25 16-24 14-24 13-20 15-20 14-18 OVERHEAD POSITION YES NO YES

COVERING/SLAG

APPROXIMATE IRON POWDER(D), %

DIGGING DEEP DIGGING DEEP

CELLULOSE/SODIUM CELLULOSE/POTASSIUM

0-10 0

MEDIUM

MEDIUM

RUTILE/SODIUM

0-10

LIGHT

LIGHT

RUTILE/POTASSIUM

0-10

LIGHT

LIGHT

RUTILE/IRON POWDER

25-40

MEDIUM

MEDIUM

0

MEDIUM

MEDIUM

MEDIUM

MEDIUM

MEDIUM

MEDIUM

LOW HYDROGEN/SODIUM LOW HYDROGEN/POTASSIUM LOW HYDROGEN/IRON POWDER IRON OXIDE/SODIUM

LIGHT

LIGHT

RUTILE/IRON POWDER

50

MEDIUM

MEDIUM

MEDIUM

MEDIUM

IRON OXIDE/IRON 50 POWDER LOW HYDROGEN/IRON 50

0 25-40 0

(A) FIRST TWO OR THREE DIGITS INDICATE TENSILE STRENGTH IN UNITS OF KSI AND OTHER MECHANICAL PROPERTIES (MECHANICAL PROPERTY REQUIREMENTS VARY WITHIN EACH CLASSIFICATION). (B) THIRD (OR) FOURTH (SECOND TO LAST) DIGIT INDICATES THE WELDING POSITION THAT CAN BE USED. (C) LAST DIGIT INDICATES USABILITY OF THE ELECTRODE. (D) IRON POWDER PERCENTAGE BASED ON WEIGHT OF THE COVERING Stainless Steel Covered Electrodes. The three-digit number that follows the prefix E indicates the chemical composition. In addition, letters or numbers can be used to indicate composition modifications or position usability. The specification AWS A5.4 identifies and classifies covered corrosion-resisting chromium and chromium-nickel steel welding electrodes. Nickel and Copper Alloys. The designations for nonferrous product classifications, such as nickel and nickel alloys in

AWS A5.11 and copper and copper alloys in AWS A5.6, follow the prefix with a list of chemical element abbreviations that are significant in identifying product composition, such as ENiCu, ENiCrFe, ECuSi, and ECuNi.

Surfacing Welding Electrodes. The designations for these products are contained in specifications AWS A5.13 and

A5.21. They are very similar to the system used to identify nonferrous electrodes. Aluminum and Aluminum Alloys. The specification for aluminum and aluminum alloy arc welding electrodes, AWS

A5.3, uses the E prefix to indicate a covered electrode, followed by a series of numbers that identify the chemical composition that is equivalent to Aluminum Association alloy designations (for example, E1100, E3003, and E4043). Suffix symbols are used in various classifications. The AWS A5.5 specification for low-alloy filler metals uses suffixes

such as -A1, -B2, -B2L, and -C1 to indicate chemical compositions. Table 3 identifies the weld deposit chemical composition associated with a number of suffixes found on low-alloy electrodes. The classifications for nonferrous products in the AWS A5.6 specification for copper alloys and in the AWS A5.11 specification for nickel alloys list a letter or number suffix that indicates position in a series of similar alloy groupings. A similar suffix pattern is also used in the AWS A5.13 and AWS A5.21 specifications for surfacing welding electrodes. Covered stainless steel electrodes employ a number, -15 or -16, as a suffix to identify usability. The -15 suffix indicates that the electrode is designed for all-position operation using DCEP electrical current. The -16 suffix indicates all-position operation with either ac or DCEP.

TABLE 3 SUFFIX SYMBOLS AND CORRESPONDING COMPOSITIONS FOR LOW-ALLOY STEEL COVERED ELECTRODES

SUFFIX(A) COMPOSITION, % C Mn A1 0.12 0.6-1.0(B) B1 0.12 0.90 B2L 0.05 0.90 B2 0.12 0.90 B3L 0.05 0.90 B3 0.12 0.90 B4L 0.05 0.90 B5 0.07-0.15 0.40-0.70 C1 0.12 1.20 C2 0.12 1.20 C3 0.12 0.40-1.25 D1 0.12 1.25-1.75 D2 0.15 1.65-2.00 G ... 1.0 MIN (C) M 0.10 0.6-2.25(B)

Si 0.40-0.80(B) 0.60-0.80(B) 0.8-1.0(B) 0.60-0.80(B) 0.8-1.0(B) 0.60-0.80(B) 1.00 0.30-0.60 0.6-0.8(B) 0.6-0.8(B) 0.80 0.6-0.8(B) 0.6-0.8(B) 0.80 MIN 0.6-0.8(B)

Ni ... ... ... ... ... ... ... ... 1.00-2.75 3.00-3.75 0.80-1.10 ... ... 0.50 MIN 1.4-2.5(B)

Cr ... 0.40-0.65 1.00-1.50 1.00-1.50 2.00-2.50 2.00-2.50 1.75-2.25 0.50-0.60 ... ... 0.15 ... ... 0.30 MIN 0.15-1.5(B)

Mo 0.40-0.65 0.40-0.65 0.40-0.65 0.40-0.65 0.90-1.20 0.90-1.20 0.40-0.65 1.00-1.25 ... ... 0.35 0.25-0.45 0.25-0.45 0.20 MIN 0.25-0.55(B)

V ... ... ... ... ... ... 0.05 ... ... 0.05 ... ... 0.10 MIN 0.05

(A) THE SUFFIX INDICATES THE CHEMICAL COMPOSITION OF THE WELD-METAL DEPOSIT. (B) AMOUNT DEPENDS ON ELECTRODE CLASSIFICATION. SINGLE VALUES INDICATE MAXIMUM, CHECK AWS A5.5 FOR THE DIFFERENT ELECTRODE CLASSES. (C) THERE ARE SEVERAL DIFFERENT M CLASSES; M CLASSIFICATIONS ARE INTENDED TO CONFORM TO MILITARY SPECIFICATIONS. Deposition Rates. The melting rate of the electrode is directly related to the welding current. The current density in the

electrode increases with higher current, which increases the melting rate, which, in turn, increases the deposition rate. The electrode coating also affects deposition rate. The iron powder types are designed to have higher deposition rates and, therefore, greater productivity. Figure 4 shows the expected deposition rate versus amperage for various electrodes at a 100% duty cycle. (The actual deposition rate will be considerably less. Deposition rate is a function of the duty cycle, which is affected by the time spent changing the electrodes, cleaning slag off of the weld, etc.) Electrode size and, therefore, usable current range are determined by the base-metal thickness, welding position, welder skill level, and joint details.

FIG. 4 DEPOSITION RATE VERSUS AMPERES FOR VARIOUS ELECTRODES

Shielded Metal Arc Welding Raymond H. Juers, Naval Surface Warfare Center

Weld Schedules and Procedures Welding schedules are tables of operating parameters that will provide high-quality welds under normal conditions.

Strict welding schedules are not as important for the manual SMAW process as they are for semiautomatic and automatic welding for several reasons. First, in a manual welding process, the welder controls conditions by arc manipulation, which achieves better control than any of the other arc welding processes. The welder also directly controls the arc voltage and travel speed and, indirectly, the welding current. Second, meter readings are rarely used in the SMAW process for the duplication of jobs. It is generally considered that the recommended welding current ranges given in Table 4 for the different types of electrodes are sufficient for most operations. The settings provide a good starting point when first welding on a new application, although they are not necessarily the only welding settings that can be used under every condition. For example, for high-production work, the current settings could be increased considerably over those shown. Factors such as weld appearance, welding position, and welder skill also allow variations from the settings.

TABLE 4 TYPICAL AMPERAGE RANGES FOR SELECTED SMAW ELECTRODES ELECTRODE DIAMETER mm in.

E6010 AND E6011

E6012

E6013

E6020

E6022

E6027 AND E7027

E7014

E7015, E7016, AND E7016-1

E7018 AND E7018-1

E70241, E7024, AND E7028

E7048

1.6

...

20-40

20-40

...

...

...

...

...

...

...

...

...

25-60

25-60

...

...

...

...

...

...

...

...

40-80

35-85

45-90

...

...

...

80140 110190 140240

80130 105180 150230

100150 130190 175250

110160 140190 170400

125185 160240 210300

65110 100150 140200 180255

70100 115165 150220 200275

100145 140190 180250 230305

...

75125 110170 140215

80125 110160 150210 200275

2.0 2.4(A) 3.2 4.0 4.8

1 16 5 64 3 32 1 8 5 32 3 16

80140 150220 210270

5.6 6.4 8.0(A)

7 32 1 4 5 16

170250 210320 275425

200320 250400 300500

210300 250350 320430

225310 275375 340450

370520 ... ...

250350 300420 375475

260340 330415 390500

240320 300390 375475

260340 315400 375470

275365 335430 400525

... ... ...

(A)

(A) THESE DIAMETERS ARE NOT MANUFACTURED IN THE E7028 CLASSIFICATION. Welder Training. The SMAW process generally requires a high degree of welder skill to consistently produce quality

welds. As a result, many training programs emphasize the SMAW process because of the arc manipulation skills developed by the welder. This acquired skill level makes the training on other processes much easier. The exact content of a training program will vary, depending on the specific application of the process. The complexity of the parts to be welded and the governing codes or specifications involved also dictate the length of the training program. For example, because a pipe welder would need more skill than a tack welder, the length of his training program would be greater. The job title of arc welder (DOT 810.384-014) describes a person whose responsibility is to weld together many components. This job includes setting up the machine and part to be welded, striking the arc and guiding it along the joint, and performing such duties as chipping, grinding, and slag removal. The welder should be able to: • • •

WELD IN ALL POSITIONS PASS EMPLOYER PERFORMANCE TESTS MEET CERTIFICATION STANDARDS OF GOVERNMENTAL AGENCIES OR PROFESSIONAL AND TECHNICAL ASSOCIATIONS

A tack welder (DOT 810.684-010) makes short beads at specific points to hold the parts in place for final welding. The tack welder also performs the duties of fitter helper. A combination welder (DOT 819-384-010) welds metal parts together to either fabricate or repair assemblies. The combination welder is able to create gas welds and electric arc welds, and to perform flame-cutting operations. The welder portion of the pipefitter description (DOT 862.381-018) must weld the pipe joint after it has been assembled and tacked in place. Procedures. There is a definite relationship between the welding current, the size of the welding electrodes, and the

welding position. These parameters must be selected to ensure that the welder has the molten weld-metal pool under complete control at all times. If the pool becomes too large, then it becomes unmanageable, perhaps allowing molten metal to run out, particularly during out-of-position welding. The welder should maintain the steady frying and crackling sound that comes with the use of correct procedures. The shape of the molten pool and the movement of the metal at the rear of the pool both serve as guides in checking weld quality. The ripples produced on the bead should be uniform and have good side tie-in with no undercut or excessive reinforcement. The factors described below are essential for maintaining high-quality welding. Correct Electrode Type. It is important to select the proper electrode for each job. The selection should be based on the type of base metal, expected service, and mechanical properties required. Correct Electrode Size. The choice of electrode size should depend on the type of electrode, welding position, joint

preparation, base-metal thickness, and welder skill. Correct Current. If the current is too high, then the electrode melts too fast and the molten pool becomes large,

irregular, and difficult to control. Current that is too low will not provide enough heat to melt the base metal, causing the molten pool to be sluggish, with a high, irregular, ropey weld bead. It should also be noted that the electrode has inherent

current limits. If the current is too high, then the core wire overheats and the coating cracks. If the current is too low, then there is insufficient heat to maintain the arc and form the protective gas shield. Correct Arc Length. If an arc is too long, then the metal melts off the electrode in large globules that wobble from side to side, resulting in a wide and irregular weld bead with considerable spatter, and, possibly, porosity and mechanical property degradation. If the arc is too short, then it has insufficient heat to melt the base metal and electrode, which often results in the electrode sticking to the work. Correct Travel Speed. A speed that is too fast allows the weld pool to freeze before impurities and gases can escape,

and the bead will be narrow and inadequate in size. When the speed is too slow, the metal piles up and the bead is larger than required. Correct Electrode Angle. In fillet welding and deep groove welding, the electrode angle is particularly important.

When making a fillet weld, the electrode should be held so that it bisects the angle between the plates and is perpendicular to the line of the weld. When undercut occurs in the vertical member, the angle should be lowered and the arc directed toward the vertical member. Correct Arc Manipulation. When weaving is used, the width of the weave and the pause at the ends of the weave become important. The welder must pause at each end of the weave to allow adequate fill buildup and fusion to occur. The welder should also quickly move across the center of the weld, because heating is more concentrated in the center than at the edges. Breaking the Arc. Before an arc is broken, it is important to know whether it will be reestablished with the next

electrode and the weld continued or whether it is the end of a weld pass. If welding is to continue, then the crater should remain and the arc quickly broken. If it is the end of a weld pass, then the arc should not be broken until the crater has filled. Correct Interbead Cleaning. Proper interbead cleaning to remove slag and any spatter is essential to the production of

high-quality welds. Proper cleaning prevents slag inclusions, "lack of fusion" defects, and porosity. Shielded Metal Arc Welding Raymond H. Juers, Naval Surface Warfare Center

Variations of the SMAW Process Gravity Welding. Gravity feed is considered to be an automatic method of applying the SMAW process. It utilizes a relatively low-cost mechanism that includes an electrode holder attached to a bracket, which slides down an inclined bar arranged along the line of weld. Special electrodes with a heavy coating are maintained in contact with the workpiece by the weight of the electrode holder and electrode. Once the process is started, it continues automatically until the electrode has burned to a short stub, whereupon the bracket and electrode holder are automatically kicked up to break the arc. One welder can operate several gravity feeders at the same time (Fig. 5). This increases productivity, reduces welder fatigue, and requires less-skilled welders, all of which result in substantial savings in welder labor costs. Table 5 compares the deposition rate, based on pounds per hour, when using one electrode manually versus two, three, four, or five gravity feeders.

TABLE 5 COMPARISON OF DEPOSITION RATES FOR CONVENTIONAL SMAW AND MULTI-ARC GRAVITY-FED WELDING Data are for making 7.9 mm (

5 in.) fillet welds using E6027 electrodes. 16

METHOD MANUAL, ONE ARC GRAVITY, TWO ARCS GRAVITY, THREE ARCS

DEPOSITION RATE, LB/H 9 17 26

GRAVITY, FOUR ARCS GRAVITY, FIVE ARCS

34 43

FIG. 5 GRAVITY FEEDERS BEING USED ON SHIP SUBASSEMBLIES

Although gravity-fed SMAW was investigated in the United States, England, and the Scandinavian countries in the late 1940s and early 1950s, credit must be given to the Japanese shipbuilders for perfecting and utilizing the process on a large scale in the early 1960s. The gravity welding process is being used in shipyards, railroad car shops, and barge yards throughout the world. It has reasonable acceptance in applications where large amounts of horizontal fillet welds must be made in a relatively small area. Firecracker welding is an automatic method by which shielded metal arc welds are made using a long electrode with an electrically nonconductive coating. Human involvement is not required after the arc is initiated. A firecracker weld is generally positioned in the flat position. The welding electrode is placed in the joint and a retaining bar is placed over it. The arc is started by shorting the end of the electrode to the workpiece. The arc length is controlled by the coating thickness. As the arc travels along the stationary electrode, the electrode melts and makes a deposit on the metal immediately underneath. Once the arc is started, the process proceeds to completion automatically. Electrodes that are up to 1 m (39 in.) long and have a core diameter of 5, 6, or 8 mm (0.20, 0.24, or 0.32 in.) have been used. Both alternating and direct current have been applied, and the former may be preferred, because of arc blow problems associated with direct current. Shielded Metal Arc Welding Raymond H. Juers, Naval Surface Warfare Center

Special Applications of the SMAW Process Underwater welding began during World War I when the British naval force used it to make temporary repairs of

leaking rivets on ship hulls. The introduction of covered electrodes enabled succesful underwater welding and the production of welds having approximately 80% of the strength and 40% of the ductility of similar welds made in air. Because of the somewhat diminished weld properties, this SMAW application is generally restricted to salvage operations or underwater repair work. Underwater welding can be subdivided into two major categories: welding in a wet environment and welding in a dry environment. Wet Welding. The relatively poor quality of welds made in a wet environment is due primarily to the problem of heat

transfer, welder visibility, and the presence of hydrogen in the arc atmosphere during the welding operation. When the base metal and the arc area are surrounded entirely by water, there is no temperature or heat buildup of the base metal at the weld, nor is preheating possible. This creates a weld-metal quench effect, which traps damaging amounts of hydrogen and also produces a weld solidification structure with reduced toughness and ductility. Both conditions contribute to the weld-metal cracking problems experienced when welding steels underwater. Another disadvantage is the restricted visibility, which is due to the equipment and the existing local contaminants in the water, as well as those generated by the welding arc. Under the most ideal conditions, welds produced in wet environments using covered electrodes are

marginal, at best. They can be placed in service for short periods of time under reduced operating conditions, but should be replaced with quality welds as quickly as possible. The covered electrodes used for wet welding must be waterproofed prior to underwater use. This can be done by wrapping them with waterproof tape or by dipping them in special sodium silicate mixes and allowing them to dry. Dry Welding. The dry environment enables the production of high-quality welds that meet all code quality requirements. The SMAW process is not very popular for welding in the dry environment, because large amounts of smoke and fumes are produced. An extensive air moving, filtering, and refrigeration system must be employed when the SMAW process is used, because a dry-environment area will quickly fill with the welding fumes, making it impossible for the welder to see that weld area and to function. For this reason, the gas-tungsten arc welding and gas-metal arc welding processes have broader use in dry welding applications.

The article "Underwater Welding" in this Volume explores the topic in greater detail. Shielded Metal Arc Welding Raymond H. Juers, Naval Surface Warfare Center

Safety Considerations The SMAW process, like all open arc welding processes, has a number of potential hazards. To alert welders about these safety concerns, all SMAW electrode containers carry a warning label that identifies the three most common safety hazards in these terms: •

•

•

ELECTRIC SHOCK CAN KILLDO NOT PERMIT ELECTRICALLY LIVE PARTS OR ELECTRODES TO CONTACT SKIN . . . OR YOUR CLOTHING OR GLOVES IF THEY ARE WETINSULATE YOURSELF FROM WORK AND GROUND FUMES AND GASES CAN BE DANGEROUS TO YOUR HEALTHKEEP FUMES AND GASES FROM YOUR BREATHING ZONE AND GENERAL AREAKEEP YOUR HEAD OUT OF FUMESUSE ENOUGH VENTILATION OR EXHAUST AT THE ARC, OR BOTH ARC RAYS CAN INJURE EYES AND BURN SKINWEAR CORRECT EYE, EAR, AND BODY PROTECTION

In addition to the general warnings on the container, further detailed information relating to safety is contained in ANSI/AWS Z49.1, "Safety in Welding and Cutting." The information described below is intended to expand on the general warnings, and should be particularly useful for SMAW welders. First, the filter plates installed within welding shields must be capable of stopping the harmful levels of infrared, ultraviolet, and visible light rays originating in the arc. Filter plates are now able to absorb 99% or more of the infrared and ultraviolet rays from the arc. The shade of the filter plate suggested for use with SMAW electrodes is given in Table 6.

TABLE 6 RECOMMENDED FILTER LENS SHADES USED IN SHIELDED METAL ARC WELDING

ELECTRODE DIAMETER LENS SHADE NO. mm in. 1 1.6 10 2.4 3.2

16 3 32 1 8

4.0 4.8 5.6 6.4 7.9 9.5

5 32 3 16 7 32 1 4 5 16 3 8

12

14

Source: ANSI/AWS Z49.1

Second, in addition to protecting himself, the welder must also be aware of others in the area who need protection, which can usually be provided by portable screens. The failure of those working around the arc to use adequate protection can result in eye burn, which is similar to sunburn and is extremely painful for a period of up to 48 h. Usually, eye burn does not permanently injure the eyes, but it can cause intense pain. A physician should be consulted in the case of severe arc burn, regardless of whether it involves the skin or the eyes. Third, the welding area must be adequately ventilated because of the heavy concentrations of smoke and fumes generated in the SMAW process. If welding is being performed in confined spaces with poor ventilation, such as in a tank, an external air supply in the form of a mask or special helmet may be required. In addition, a second person should be stationed at the tank manhole to provide any necessary assistance. Special ventilation is required when welding stainless steels or metals coated with copper, zinc, lead, or cadmium, because of the toxic nature of these fumes. Fourth, cables with frayed or cracked insulation and faulty or badly worn connections can cause electrical short circuits and electrical shocks to personnel. When it is necessary to weld in a damp or wet area, the welder should wear rubber boots and stand on a dry, insulated platform. Shielded Metal Arc Welding Raymond H. Juers, Naval Surface Warfare Center

Selected References

• H.B. CARY, MODERN WELDING TECHNOLOGY, 2ND ED., PRENTICE-HALL, 1989 • "FILLER METAL COMPARISON CHARTS," AWS FMC-89, AMERICAN WELDING SOCIETY • "SAFETY IN WELDING AND CUTTING," ANSI/AWS Z49.1, AMERICAN WELDING SOCIETY • "TECHNICAL GUIDE FOR SHIELDED METAL ARC WELDING," HOBART BROTHERS CO. Gas-Metal Arc Welding D.B. Holliday, Westinghouse Electric Corporation

Introduction GAS-METAL ARC WELDING (GMAW) is an arc welding process that joins metals together by heating them with an electric arc that is established between a consumable electrode (wire) and the workpiece. An externally supplied gas or gas mixture acts to shield the arc and molten weld pool. Although the basic GMAW concept was introduced in the 1920s, it was not commercially available until 1948. At first, it was considered to be fundamentally a high-current-density, small-diameter, bare-metal electrode process using an inert

gas for arc shielding. Its primary application was aluminum welding. As a result, it became known as metal-inert gas (MIG) welding, which is still common nomenclature. Subsequent process developments included operation at low current densities and pulsed direct current, application to a broader range of materials, and the use of reactive gases (particularly carbon dioxide) and gas mixtures. The latter development, in which both inert and reactive gases are used, led to the formal acceptance of the term gas-metal arc welding. The GMAW process can be operated in semi-automatic and automatic modes. All commercially important metals, such as carbon steel, high-strength low-alloy steel, stainless steel, aluminum, copper, and nickel alloys can be welded in all positions by this process if appropriate shielding gases, electrodes, and welding parameters are chosen. Advantages. The applications of the process are dictated by its advantages, the most important of which are: • • • • • •

•

•

ELECTRODE LENGTH DOES NOT FACE THE RESTRICTIONS ENCOUNTERED WITH SHIELDED-METAL ARC WELDING (SMAW). WELDING CAN BE ACCOMPLISHED IN ALL POSITIONS, WHEN THE PROPER PARAMETERS ARE USED, A FEATURE NOT FOUND IN SUBMERGED ARC WELDING. WELDING SPEEDS ARE HIGHER THAN THOSE OF THE SMAW PROCESS. DEPOSITION RATES ARE SIGNIFICANTLY HIGHER THAN THOSE OBTAINED BY THE SMAW PROCESS. CONTINUOUS WIRE FEED ENABLES LONG WELDS TO BE DEPOSITED WITHOUT STOPS AND STARTS. PENETRATION THAT IS DEEPER THAN THAT OF THE SMAW PROCESS IS POSSIBLE, WHICH MAY PERMIT THE USE OF SMALLER-SIZED FILLET WELDS FOR EQUIVALENT STRENGTHS. LESS OPERATOR SKILL IS REQUIRED THAN FOR OTHER CONVENTIONAL PROCESSES, BECAUSE THE ARC LENGTH IS MAINTAINED CONSTANT WITH REASONABLE VARIATIONS IN THE DISTANCE BETWEEN THE CONTACT TIP AND THE WORKPIECE. MINIMAL POSTWELD CLEANING IS REQUIRED BECAUSE OF THE ABSENCE OF A HEAVY SLAG.

These advantages make the process particularly well suited to high-production and automated welding applications. With the advent of robotics, gas-metal arc welding has become the predominant process choice. Limitations. The GMAW process, like any welding process, has certain limitations that restrict its use: • •

THE WELDING EQUIPMENT IS MORE COMPLEX, USUALLY MORE COSTLY, AND LESS PORTABLE THAN SMAW EQUIPMENT. THE PROCESS IS MORE DIFFICULT TO APPLY IN HARD-TO-REACH PLACES BECAUSE THE WELDING GUN IS LARGER THAN A SMAW HOLDER AND MUST BE HELD CLOSE TO THE JOINT (WITHIN 10 TO 19 MM, OR

•

•

3 8

TO

3 4

IN.) TO ENSURE THAT THE WELD METAL IS

PROPERLY SHIELDED. THE WELDING ARC MUST BE PROTECTED AGAINST AIR DRAFTS THAT CAN DISPERSE THE SHIELDING GAS, WHICH LIMITS OUTDOOR APPLICATIONS UNLESS PROTECTIVE SHIELDS ARE PLACED AROUND THE WELDING AREA. RELATIVELY HIGH LEVELS OF RADIATED HEAT AND ARC INTENSITY CAN HINDER OPERATOR ACCEPTANCE OF THE PROCESS.

Gas-Metal Arc Welding D.B. Holliday, Westinghouse Electric Corporation

Process Fundamentals Principles of Operation. In the GMAW process (Fig. 1), an arc is established between a continuously fed electrode of filler metal and the workpiece. After proper settings are made by the operator, the arc length is maintained at the set value, despite the reasonable changes that would be expected in the gun-to-work distance during normal operation. This automatic arc regulation is achieved in one of two ways. The most common method is to utilize a constant-speed (but adjustable) electrode feed unit with a variable-current (constant-voltage) power source. As the gun-to-work relationship changes, which instantaneously alters the arc length, the power source delivers either more current (if the arc length is decreased) or less current (if the arc length is increased). This change in current will cause a corresponding change in the electrode melt-off rate, thus maintaining the desired arc length.

FIG. 1 SCHEMATIC OF GMAW PROCESS

The second method of arc regulation utilizes a constant-current power source and a variable-speed, voltage-sensing electrode feeder. In this case, as the arc length changes, there is a corresponding change in the voltage across the arc. As this voltage change is detected, the speed of the electrode feed unit will change to provide either more or less electrode per unit of time. This method of regulation is usually limited to larger electrodes with lower feed speeds. Metal Transfer Mechanisms. The characteristics of the GMAW process are best described by reviewing the three basic means by which metal is transferred from the electrode to the work: short-circuiting transfer, globular transfer, or spray transfer. The type of transfer is determined by a number of factors, the most influential of which are: • • • • • •

MAGNITUDE AND TYPE OF WELDING CURRENT ELECTRODE DIAMETER ELECTRODE COMPOSITION ELECTRODE EXTENSION BEYOND THE CONTACT TIP OR TUBE SHIELDING GAS POWER SUPPLY OUTPUT

Short-circuiting transfer encompasses the lowest range of welding currents and electrode diameters associated with the

GMAW process. This type of transfer produces a small, fast-freezing weld pool that is generally suited for joining thin sections, for out-of-position welding, and for bridging of large root openings. Metal is transferred from the electrode to the workpiece only during a period when the electrode is in contact with the weld pool, and there is no metal transfer across the arc gap (Fig. 2).

FIG. 2 TRANSFER MODES IN GMAW PROCESS

The electrode contacts the molten weld pool at a steady rate that can range from 20 to over 200 times per second. As the wire touches the weld metal, the current increases and the liquid metal at the wire tip is pinched off, initiating an arc. The rate of current increase must be high enough to heat the electrode and promote metal transfer, yet low enough to minimize spatter caused by violent separation of the molten drop. The rate of current increase is controlled by adjusting the power source inductance. The optimum setting depends on the electrical resistance of the welding circuit and the melting temperature of the electrode. When the arc is initiated, the wire melts at the tip as it is fed forward toward the next short circuit. The open-circuit voltage of the power source must be low enough so that the drop of molten metal cannot transfer until it contacts the weld metal. Because metal transfer only occurs during short circuiting, the shielding gas has very little effect on the transfer itself. However, the gas does influence the operating characteristics of the arc and the base-metal penetration. The use of carbon dioxide generally produces high spatter levels, when compared with inert gases, but it allows deeper penetration when welding steels. To achieve a good compromise between spatter and penetration, mixtures of carbon dioxide and argon are often used. With nonferrous metals, argon-helium mixtures are used to achieve this compromise. Globular Transfer. With a positive electrode, globular transfer takes place when the current density is relatively low,

regardless of the type of shielding gas. However, the use of carbon dioxide or helium results in this type of transfer at all usable welding currents. Globular transfer is characterized by a drop size with a diameter that is greater than that of the electrode. This large drop is easily acted upon by gravity, which limits successful transfer to the flat position.

At average currents that are slightly higher than those used in short-circuiting transfer, axially directed globular transfer can be achieved in a substantially inert gas shield. However, if the arc length is too short, then the enlarging drop can short to the workpiece, become superheated, and disintegrate, producing considerable spatter. Therefore, the arc length must be long enough to ensure that the drop detaches before it contacts the weld pool. However, when higher voltage values are used, the weld is likely to be unacceptable, because of a lack of fusion, insufficient penetration, and excessive reinforcement. This limits the use of this transfer mode to very few production applications. Carbon dioxide shielding produces a randomly directed globular transfer when the welding current and voltage values are significantly higher than the range used for short-circuiting transfer. Although severe spatter conditions result when conventional techniques are used, carbon dioxide is still the most commonly used shielding gas for welding mild steel when the quality requirements are not too rigorous. The spatter problem is controlled by "burying" the arc below the weld/base-metal surface. The resulting arc forces are adequate enough to produce a depression that contains the spatter. This technique requires relatively high currents and results in very deep penetration. Good operator setup skills are required. However, poor wetting action can result in an excessive weld reinforcement. Spray Transfer. A very stable, spatter-free "spray" transfer mode can be produced when argon-rich shielding is used. This type of transfer requires the use of direct current with the electrode positive and a current level that is above a critical value called the "transition current." Below this current level, transfer occurs in the globular mode at the rate of a few drops per second. At values above the transition current, transfer occurs in the form of very small drops that are formed and detached at the rate of hundreds per second and are accelerated axially across the arc gap.

The transition current is proportional to the electrode diameter, and, to a lesser extent, to the electrode extension. It also has a direct relationship to the filler metal melting temperature. Transition currents for various materials and electrode diameters are shown in Table 1.

TABLE 1 GLOBULAR-TO-SPRAY TRANSITION CURRENTS FOR SELECTED ELECTRODES

WIRE ELECTRODE TYPE LOW-CARBON STEEL

STAINLESS STEEL

SHIELDING WIRE DIAMETER GAS mm in. 98AR-2O2 0.58 0.023 0.76 0.030 0.89 0.035 1.14 0.045 1.57 0.062 95AR-5O2 0.89 0.035 1.14 0.045 1.57 0.062 92AR-8CO2 0.89 0.035 1.14 0.045 1.57 0.062 85AR-15CO2 0.89 0.035 1.14 0.045 1.57 0.062 80AR-20CO2 0.89 0.035 1.14 0.045 1.57 0.062 99AR-1O2 0.89 0.035 1.14 0.045 1.57 0.062 AR-HE-CO2 0.89 0.035 1.14 0.045 1.57 0.062 AR-H2-CO2 0.89 0.035

SPRAY ARC CURRENT, A 135 150 165 220 275 155 200 265 175 225 290 180 240 295 195 255 345 150 195 265 160 205 280 145

ALUMINUM

ARGON

DEOXIDIZED COPPER ARGON

SILICON BRONZE

ARGON

1.14 1.57 0.76 1.19 1.57 0.89 1.14 1.57 0.89 1.14 1.57

0.045 0.062 0.030 0.047 0.062 0.035 0.045 0.062 0.035 0.045 0.062

185 255 95 135 180 180 210 310 165 205 270

Source: Union Carbide Industrial Gases

The spray transfer mode results in a highly directed stream of discrete drops that are accelerated by arc forces to velocities that overcome the effects of gravity. This enables the process to be used in any position, under certain conditions. Because the drops are separated, short circuits do not occur, and the spatter level is negligible, if not totally eliminated. Another characteristic of spray transfer is the "finger" penetration pattern that it produces directly below the electrode tip. Although the penetration can be deep, it can be affected by magnetic fields that must be controlled to ensure that it is always located at the center of the weld penetration profile. Otherwise, a lack of fusion and an irregular bead surface profile can result. The spray transfer mode can be used to weld almost any metal or alloy, because of the inert characteristics of the argon shield. Sometimes, thickness can be a factor, because of the relatively high current levels required. The resultant arc forces can cut through, rather than weld, thin sheets. In addition, high deposition rates can result in a weld pool size that cannot be supported by surface tension in the vertical and overhead positions. However, the thickness and position limitations of spray transfer have been largely overcome by specially designed power supplies. These machines produce carefully controlled current outputs that "pulse" the welding current from levels below the transition current to levels above it. Figure 3 shows the two levels of current provided by these machines. One is a constant, low-background current that sustains the arc without providing enough energy to cause the formation of drops on the wire tip. The other is a superimposed pulsing current with an amplitude that is greater than the transition current necessary for spray transfer. During this pulse, one or more drops are formed and transferred. The frequency and amplitude of the pulses control the energy level of the arc and, therefore, the rate at which the wire melts. By reducing the arc energy and the wire melting rate, it is possible to retain many of the desirable features of spray transfer while joining sheet metals and welding thick metals in all positions.

FIG. 3 CHARACTERISTIC CURRENT WAVEFORM FOR A "PULSING" POWER SUPPLY

Many variations of such machines are available. The simplest provide a single frequency of pulsing (60 or 120 pulses/s) and independent control of the background and pulsing current levels. Synergic machines, which are sophisticated, automatically provide the optimum combination of background and pulsing current levels for any given setting of wire feed speed. Normally, these settings are specific to an electrode/shielding gas combination and must be changed or reprogrammed when the combination is changed. Process Variables. The important variables of the GMAW process that affect weld penetration, bead geometry, and

overall weld quality are: • • • • • • •

WELDING CURRENT (ELECTRODE FEED SPEED) POLARITY ARC VOLTAGE (ARC LENGTH) TRAVEL SPEED ELECTRODE EXTENSION ELECTRODE ORIENTATION (GUN ANGLE) ELECTRODE DIAMETER

Knowledge and control of these variables are essential to consistently produce welds of satisfactory quality. Because they are not completely independent of one another, changing one variable generally requires changing one or more of the others to produce the desired results. The effects of these variables on deposit attributes are shown in Table 2.

TABLE 2 EFFECT OF CHANGES IN PROCESS VARIABLES ON WELD ATTRIBUTES WELDING VARIABLES TO CHANGE

CURRENT AND WIRE FEED SPEED VOLTAGE TRAVEL SPEED ELECTRODE EXTENSION WIRE DIAMETER SHIELD GAS % GUN ANGLE

DESIRED CHANGES PENETRATION DEPOSITION RATE

BEAD SIZE

INCREASE

DECREASE

INCREASE

DECREASE

INCREASE

DECREASE

INCREASE

DECREASE

INCREASE

DECREASE

INCREASE

DECREASE

INCREASE

DECREASE

LITTLE EFFECT

LITTLE EFFECT

NO EFFECT NO EFFECT DECREASE

NO EFFECT NO EFFECT INCREASE

LITTLE EFFECT LITTLE EFFECT INCREASE(A)

LITTLE EFFECT LITTLE EFFECT DECREASE(A)

LITTLE EFFECT DECREASE

LITTLE EFFECT INCREASE

INCREASE

DECREASE

DECREASE

INCREASE

INCREASE

DECREASE

DECREASE

INCREASE

DECREASE

INCREASE

DECREASE

INCREASE

INCREASE

DECREASE

LITTLE EFFECT DECREASE

PUSH

LITTLE EFFECT LITTLE EFFECT

LITTLE EFFECT LITTLE EFFECT LITTLE EFFECT

LITTLE EFFECT INCREASE

DRAG

LITTLE EFFECT LITTLE EFFECT

LITTLE EFFECT LITTLE EFFECT LITTLE EFFECT

PUSH

DRAG

BEAD WIDTH

(A) WILL RESULT IN DESIRED CHANGE IF CURRENT LEVELS ARE MAINTAINED BY ADJUSTMENT OF WIRE FEED SPEED Considerable skill and experience are necessary to select the optimal combination for each application. This selection is further complicated by the fact that the optimal settings for the variables are also affected by the type of base metal, the electrode composition, the welding position, quality requirements, and the number of completed weldments required. Thus, no single set of parameters provides optimal results in every case. Welding Current. As the electrode feed speed is varied, the welding current varies in a like manner when a constantvoltage power source is used. This occurs because the current output of the power source varies dramatically with the slight changes in the arc voltage (arc length) that result when changes are made in the electrode feed speed. When all

other variables are held constant, an increase in welding current results in an increase in the depth and width of penetration, deposition rate, and weld bead size. Polarity is the term used to described the electrical connection of the welding gun in relation to the terminals of a direct-

current (dc) power source. When the gun power lead is connected to the positive terminal, the polarity is designated as direct current, electrode positive (DCEP). Alternatively, a connection to the negative terminal is designated as direct current, electrode negative (DCEN). The vast majority of GMAW applications utilize DCEP, because it provides for a stable arc, low spatter, a good weld bead profile, and the greatest depth of penetration. Arc voltage and arc length are related terms that are often used interchangeably. However, they are different. Arc voltage

is an approximate means of stating the physical arc length in electrical terms. The same physical arc length, however, could yield different arc voltage readings, depending on factors such as shielding gas, current, and electrode extension. When all variables are held constant, a reliable relationship exists between the two: an increase in voltage setting will result in longer arc length. Although the arc length is the variable of interest and the one that should be controlled, arc voltage is more easily monitored. Because of this fact, and because the arc voltage is normally required to be specified in welding procedures, it is the term that is more commonly used. From any specific value of arc voltage, an increase tends to flatten the weld bead and increase the width of the fusion zone. Excessively high voltage can cause porosity, spatter, and undercut. A reduction in voltage results in a narrower weld bead with a higher crown. Travel speed is the linear rate at which the arc is moved along the weld joint. When all other conditions are held

constant, weld penetration is a maximum at an intermediate travel speed. When travel speed is decreased, the filler metal deposition per unit length increases. At very slow speeds, the welding arc impinges on the molten weld pool, rather than the base metal, thereby reducing the effective penetration. As the travel speed is increased, the thermal energy transmitted to the base metal from the arc increases, because the arc acts more directly on the base metal. However, further increases in travel speed impart less thermal energy to the base metal. Thus, melting of the base metal first increases and then decreases with increasing travel speed. As travel speed is increased further, there is a tendency toward undercutting along the edges of the weld bead, because there is insufficient deposition of filler metal to fill the path melted by the arc. Electrode orientation is described in two ways: by the relationship of the electrode axis with respect to the direction of

travel (the travel angle) and by the angle between the electrode axis and the adjacent work surface (work angle). When the electrode points in a direction opposite to the travel direction, it results in a trail angle and is known as the backhand welding technique. When the electrode points in the direction of travel, it results in a lead angle and is called the forehand welding technique. For all positions, a trailing travel angle that ranges from 5 to 15° (from perpendicular) provides a weld with maximum penetration and a narrow, convex surface configuration. It also provides for maximum shielding of the molten weld pool. However, the common technique utilizes a leading travel angle, which provides better visibility for the operator and a weld with a flatter surface profile. For some materials, such as aluminum, a leading angle is preferred, because it provides a "cleaning action" ahead of the molten weld metal, which promotes wetting and reduces base-material oxidation. When producing fillet welds in the horizontal position, the work angle should be about 45° to the vertical member. The electrode extension is the distance between the last point of electrical contact (usually the gun contact tip or tube)

and the end of the electrode. An increase in the amount of this extension causes an increase in electrical resistance. This, in turn, generates additional heat in the electrode, which contributes to greater electrode melting rates. Without an increase in arc voltage, the additional metal will be deposited as a narrow, high-crowned weld bead. The optimum electrode extension generally ranges from 6.4 to 13 mm ( (

1 1 to in.) for short-circuiting transfer and from 13 to 25 mm 4 2

1 to 1 in.) for spray and globular transfers. 2

The electrode diameter influences the weld bead configuration. A larger electrode requires a higher minimum current

than a smaller electrode does to achieve the same metal transfer characteristics.

Higher currents, in turn, produce additional electrode melting and larger, more-fluid weld deposits. Higher currents also result in higher deposition rates and greater penetration, but may prevent the use of some electrodes in the vertical and overhead positions. Gas-Metal Arc Welding D.B. Holliday, Westinghouse Electric Corporation

Equipment The basic equipment for a typical GMAW installation is shown in Fig. 1. The major components are discussed below. A welding gun provides electrical current to the electrode, directs it to the workpiece, and provides a vehicle for

directing shielding gas to the weld area. Different types of guns have been designed for many varied applications, ranging from heavy-duty guns for high-current, high-volume production to lightweight guns for low-current or out-of-position welding. The most commonly used guns are designed to be cooled by the surrounding air (Fig. 4). However, as amperage requirements increase, a water-cooled gun may be required. Guns are rated based on their current-carrying capacity, generally with a CO2 shielding gas. If inert gases are used, these gun ratings must be reduced significantly. Guns can also be equipped with their own integral electrode feed units.

FIG. 4 CROSS SECTION OF AN AIR-COOLED GMAW GUN

The contact tube, usually made of copper or a copper alloy, is used to transmit welding current to the electrode, as well as to direct the electrode toward the work. The contact tube is connected electrically to the welding power supply by the power cable. The inner surface of the contact tube is very important, because the electrode must feed easily through this tube while making a good electrical contact. Generally, the hole in the contact tube would be from 0.13 to 0.25 mm (0.005 to 0.010 in.) larger than the wire being used, although larger sizes may be required for materials such as aluminum. The hole should be checked periodically and replaced if it has become elongated because of excessive wear. If a tip in this condition is used, it can result in poor electrical contact and erratic arc characteristics. The nozzle directs an even-flowing column of shielding gas into the welding zone. It is extremely important to maintain an even flow in order to adequately protect the molten weld metal from atmospheric contamination. Different-sized nozzles are available and should be chosen according to the application, that is, larger nozzles for high-current work where the puddle is large and smaller nozzles for low-current and short-circuiting welding. The electrode conduit and liner are connected to a bracket adjacent to the feed rolls on the electrode feed motor. The conduit and liner support both protect and direct the electrode from the feed rolls to the gun and contact tube. Uninterrupted electrode feeding is necessary to ensure good arc stability. A steel liner is recommended when using hard

electrode materials, such as steel and copper, whereas nylon liners should be used for soft electrode materials, such as aluminum and magnesium. The electrode feed unit, or wire feeder, consists of an electric motor, output shaft, drive rolls, and accessories for

maintaining electrode alignment and pressure (Fig. 5). These units can be separate or integrated with the speed control or located remotely from it. The electrode feed motor is usually a direct-current type and provides the mechanical energy for pushing the electrode through the gun and to the work. It has a control circuit that varies the motor speed over a broad range. The feed unit can be an integral component of the gun (Fig. 6) or dual-feed units, one in the gun and one in a separate feeder, can be electrically coupled together to provide a "push-pull" system. The spool-gun and push-pull systems are often used on aluminum, where difficulty can be encountered in trying to push the wire through a conduit to the gun.

FIG. 5 TYPICAL REMOTE-WIRE FEED UNITS. (A) SCHEMATIC OF A TYPICAL FOUR-ROLL SYSTEM MECHANISM. (B) CLOSE-UP VIEW OF A TYPICAL TWO-ROLL SYSTEM GEARBOX

FIG. 6 SPOOL-ON-GUN TYPE TORCH WITH COVER REMOVED TO SHOW SELF-CONTAINED PULL-TYPE FEEDUNIT SYSTEM

The feed motor is connected through a gear reducer to a set of wire feed rolls that transmit mechanical energy to the electrode, pulling it from the source and pushing it through the welding gun. Various types of rolls are available, including knurled, "U" groove, "V" groove, and flat. The knurled design is used for harder wires, such as steel, and allows maximum frictional force to be transmitted to the wire with a minimum of drive roll pressure. These types of rolls are not recommended for soft wires, such as aluminum, because they tend to cause the wire to flake, which can eventually clog the gun or liner. For these softer wires, the "U" groove or "V" groove type will allow the application of uniform pressure around the wire without deforming it. The flat rolls can be used with the smaller-diameter wires and in combination with a "U" or "V" groove. The welding control mechanism and the electrode feed motor for semiautomatic operation are usually provided in one

integrated package. The main function of the welding control mechanism is to regulate the speed of the electrode feed motor, usually through an electronic governor. The control also regulates the starting and stopping of the electrode feed through a signal received from the gun switch.

Normally, shielding gas, water (when used), and welding power are also delivered to the gun through the control mechanism, which requires direct connection to these facilities and the power supply. Gas and water flow are regulated to coincide with the weld start and stop by using solenoid valves. The control mechanism can also sequence gas flow starts and stops and energize the power source contractor. The control mechanism may allow some gas to flow before welding starts (preflow) and after welding stops (postflow) to protect the molten weld puddle. The control mechanism is usually independently powered by 115 V ac. The welding power source provides suitable electrical power (generally 20 to 80 V) that is delivered to the electrode and workpiece to produce the arc. Because the vast majority of GMAW applications utilize DCEP, the positive lead is connected to the gun and the negative lead, to the workpiece. The power source can be the "static" type in which incoming utility power (120 to 480 V) is reduced to welding voltage by a transformer or solid-state inverter. It could also be the "rotating" type in which the welding power is provided by a rotating generator driven by a motor or internal combustion engine. The static type is normally used in shops where there is an available source of power. It has advantages over the rotating type in that it can respond more rapidly to varying arc conditions. The rotating type is generally used at field sites where external power in unavailable. Both types of power sources can be designed and built to provide either a constant current or constant potential (cp) output, the latter of which is the most common by far. On newer power sources, this cp output can be pulsed at either a constant or variable frequency. With the advent of solidstate electronic power sources, such as inverters, even further control over the pulsing variables (for example, frequency, pulse width, and so on) can be obtained.

When used in conjunction with a constant-speed wire feeder, the constant-voltage power source compensates for the variations in the contact-tip-work-distance that can occur during normal welding operations. It does this by instantaneously increasing or decreasing welding current to increase or decrease the electrode burnoff rate. The initial arc length is established by adjusting the voltage at the power source. Once this is set, no other changes are required during welding. The wire feed speed, which is also the current control, is then set by the operator and adjusted as necessary. In addition to this self-regulating feature of the CP power source, control over slope and inductance is included on those machines intended for short-circuiting transfer. Additional controls are also provided when using power sources that have pulsing capabilities. Additional information is available in the article "Power Sources" in this Volume. Electrode Source. The GMAW process uses a continuously fed electrode that is consumed at relatively high speeds.

Therefore, the electrode source must provide a large volume of material that can readily be fed to the gun to ensure maximum process efficiency. This source is usually in the form of a spool or coil that can hold from 7 to 27 kg (15 to 60 lb) of wire that has been wound to allow free feeding without kinks or tangles. Larger spools of up to 115 kg (250 lb) are also available, and material can be provided in drums of 340 to 455 kg (750 to 1000 lb). Small spools of 0.45 to 0.9 kg (1 to 2 lb) are used for spool-on-gun equipment. Regulated Shielding Gas Supply. A system is required to provide constant shielding gas pressure and flow rate during

welding. This system consists of a regulator connected to a supply of "welding grade" shielding gas, as well as the necessary hoses or piping. The regulator is a device that reduces the source gas pressure to a constant working pressure, regardless of variations at the source. It can be a single-stage or dual-stage type and may have a built-in flowmeter. The shielding gas source can be a high-pressure cylinder, a liquid-filled cylinder, or a bulk-liquid tank. Gas mixtures are available in a single cylinder. Mixing devices can also be used to obtain the correct proportions when two or more gases or liquids are used. The type and size of the gas storage source depend on economic considerations that are based on the volume of shielding gas consumed per unit of time. Gas-Metal Arc Welding D.B. Holliday, Westinghouse Electric Corporation

Consumables The two consumable, but essential, elements of the GMAW process are the electrode and the shielding gas, each of which is described below. The chemical composition of the electrode must be selected to achieve the desired properties in the weld metal. The composition is designed with extra deoxidizers or other scavenging agents to compensate for reactions with the

atmosphere and the base metal. The deoxidizers most commonly used in steel electrodes are silicon and manganese. Silicon can also be added in all transfer modes to increase weld metal fluidity or it can be added when a 300 series stainless steel electrode is used. The physical characteristics (finish, straightness, and others) of electrodes used in the GMAW process are important

to successful welding. The material specifications for these electrodes establish manufacturing requirements to ensure that users receive a uniform product that feeds smoothly through the equipment and has these characteristics, as well: • • • •

UNIFORM WINDING ON THE SPOOL OR COIL WITH NO KINKS OR BENDS SMOOTH SURFACE FINISH FREE OF SLIVERS, SCRATCHES, OR SCALE PRESCRIBED CAST AND HELIX UNIFORM DIAMETER

Cast and helix refer to dimensions of a single coil of wire removed from a spool or coil and layed (that is, cast) on a flat surface. If this coil is too small in diameter (cast) or shows an excessive lift from the flat surface (helix) wire, feeding problems during welding can be anticipated. Shielding Gas The primary function of the shielding gas in most of the welding processes is to protect the surrounding atmosphere from contact with molten metal. In the GMAW process, this gas plays an additional role in that it has a pronounced effect on arc characteristics, mode of metal transfer, depth of fusion, weld bead profile, welding speed, and cleaning action. Inert gases, such as argon and helium, are commonly used, as is the active gas, CO2. It is also common to use mixtures of these gases and to employ small additions of oxygen. Information about shielding gas compositions and about which gases to use for specific joining applications is provided in the article "Shielding Gases" in this Volume. Gas-Metal Arc Welding D.B. Holliday, Westinghouse Electric Corporation

Safety The major hazards of concern during GMAW are: the fumes and gases, which can harm health; the high-voltage electricity, which can injure and kill; the arc rays, which can injure eyes and burn skin; and the noise which may be present that can damage hearing. The type and amount of fumes and gas present during welding depend on the electrode being used, the alloy being welded, and the presence of any coatings on the base metal. To guard against potential hazards, a welder should keep his head out of the fume plume and avoid breathing the fumes and gases caused by the arc. Ventilation is always required. Electrode shock can result from exposure to the high open-circuit voltages associated with welding power supplies. All electrical equipment and the workpiece must be connected to an approved electrical ground. Cables should be of sufficient size to carry the maximum current required. Insulation should be protected from cuts and abrasion, and the cable should not come into contact with oils, paints, or other fluids which may cause deterioration. Work areas, equipment, and clothing must be kept dry at all times. The welder should be well insulated, wearing dry gloves, rubbersoled shoes, and standing on a dry board or platform while welding. Radiant energy, especially in the ultraviolet range, is intense during GMAW. To protect the eyes from injury, the proper filter shade for the welding-current level selected should be used. These greater intensities of ultraviolet radiation can cause rapid disintegration of cotton clothing. Leather, wool, and aluminum-coated cloth will better withstand exposure to arc radiation and better protect exposed skin surfaces.

When noise has been determined to be excessive in the work area, ear protection should be used. This can also be used to prevent spatter from entering the ear. Conventional fire prevention requirements, such as removal of combustibles from the work area, should be followed. Sparks, slag, and spatter can travel long distances, so care must be taken to minimize the start of a fire at locations removed from the welding operation. For further information, see the guidelines set forth in the National Fire Protection Association Standard NFPA No. 51B, "Fire Protection in Use of Cutting and Welding Processes." Care should be exercised in the handling, storage, and use of cylinders containing high-pressure and liquefied gases. Cylinders should be secured by chains or straps during handling or use. Approved pressure-reducing regulators should be used to provide a constant, controllable working pressure for the equipment in use. Lubricants or pipe fitting compounds should not be used for making any connections, as they can interfere with the regulating equipment, and in the case of oxygen service, they can contribute to a catastrophic fire. For further safety information, see ANSI/AWS Z49.1, "Safety in Welding and Cutting." Flux-Cored Arc Welding David W. Meyer, The Esab Group, Inc.

Introduction IN THE FLUX-CORED ARC WELDING (FCAW) process, the heat for welding is produced by an electric arc between a continuous filler metal electrode and the workpiece. A tubular, flux-cored electrode makes this welding process unique. The flux contained within the electrode can make the electrode self-shielding. Alternatively, an external shielding gas may be required. Flux-Cored Arc Welding David W. Meyer, The Esab Group, Inc.

Process Features Flux-cored arc welding has two major variations. The gas-shielded FCAW process (Fig. 1) uses an externally supplied gas to assist in shielding the arc from nitrogen and oxygen in the atmosphere. Generally, the core ingredients in gasshielded electrodes are slag formers, deoxidizers, arc stabilizers, and alloying elements.

FIG. 1 GAS-SHIELDED FLUX-CORED ARC WELDING. SOURCE: REF 1

In the self-shielded FCAW process, the core ingredients protect the weld metal from the atmosphere without external shielding. Some self-shielded electrodes provide their own shielding gas through the decomposition of core ingredients. Others rely on slag shielding, where the metal drops being transferred across the arc and the molten weld pool are protected from the atmosphere by a slag covering. Many self-shielded electrodes also contain substantial amounts of deoxidizing and denitrifying ingredients to help achieve sound weld metal. Self-shielded electrodes can also contain arc stabilizers and alloying elements. Advantages. Because it combines the productivity of continuous welding with the benefits of having a flux present, the

FCAW process has several advantages relative to other welding processes. These advantages include: • • • • •

HIGH DEPOSITION RATES, ESPECIALLY FOR OUT-OF-POSITION WELDING LESS OPERATOR SKILL REQUIRED THAN FOR GAS-METAL ARC WELDING (GMAW) SIMPLER AND MORE ADAPTABLE THAN SUBMERGED ARC WELDING (SAW) DEEPER PENETRATION THAN SHIELDED METAL ARC WELDING (SMAW) MORE TOLERANT OF RUST AND MILL SCALE THAN GMAW

Disadvantages of the FCAW process include: • •

SLAG MUST BE REMOVED FROM THE WELD AND DISPOSED OF MORE SMOKE AND FUME ARE PRODUCED IN FCAW THAN IN THE GMAW AND SAW PROCESSES

• •

FUME EXTRACTION IS GENERALLY REQUIRED EQUIPMENT IS MORE COMPLEX AND MUCH LESS PORTABLE THAN SMAW EQUIPMENT


1. WELDING HANDBOOK, 8TH ED., VOL 2, AWS, 1991 Flux-Cored Arc Welding David W. Meyer, The Esab Group, Inc.

Applications Flux-cored arc welding enjoys widespread use in many industries. Both the gas-shielded and self-shielded FCAW processes are used to fabricate structures from carbon and low-alloy steels. Both process variants are used for shop fabrication, but the self-shielded FCAW process is preferred for field use. The acceptability of the FCAW process for structural use is illustrated by the fact that prequalified joints are included in the structural welding code of the American Welding Society (AWS). Gas-shielded flux-cored electrodes are commonly used to weld carbon, low-alloy steel, and stainless steels in the construction of pressure vessels and piping for the chemical processing, petroleum refining, and power-generation industries. In addition, flux-cored electrodes are used to weld some nickel-base alloys. Flux-cored electrodes are also used in the automotive and heavy-equipment industries in the fabrication of frame members, axle housings, wheel rims, suspension components, and other parts. Small-diameter flux-cored electrodes are used for automotive body repair. Flux-Cored Arc Welding David W. Meyer, The Esab Group, Inc.

Equipment The FCAW process utilizes semiautomatic, mechanized, and fully automatic welding systems. The basic equipment includes a power supply, wire feed system, and welding gun. The required auxiliary equipment, such as shielding gas, depends on the process variant used and the degree of automation. Fume removal equipment must also be considered in most applications of the FCAW process. Typical semiautomatic equipment is shown in Fig. 2. The equipment used in the gas-shielded FCAW process is typically identical to GMAW equipment.

FIG. 2 SEMIAUTOMATIC FCAW EQUIPMENT. SOURCE: REF 1

The recommended power supply for the semi-automatic FCAW process is a constant-voltage direct current (dc) machine. Most power supplies used for semiautomatic FCAW have output ratings of 600 A or less. A power supply rated at 60% or more duty cycle is the best choice for most industrial applications, whereas a duty-cycle rating as low as 20% may be sufficient for maintenance and repair applications. Constant-current power supplies are used in certain situations, such as field welding applications, where portable constant-current SMAW power supplies are readily available. The addition of a contactor and a voltage-sensing wire feeder makes this an adequate welding system. However, such a system is only recommended when the use of a constantvoltage system is not feasible, because constant-current systems produce an inherently less-stable welding arc than constant-voltage systems. Wire feeders for constant-voltage FCAW systems are generally simple devices that provide a constant wire feed speed. The power supply provides sufficient current to maintain an arc at the voltage that is preset at the power supply. A change in wire feed speed results in a change in the welding current. In a constant-current system, the wire feeder is somewhat more complex. The welding current is preset at the power supply. The wire feeder has a voltage-sensing feedback loop that allows it to adjust the wire feed speed to maintain the desired welding voltage. The wire feeder generally contains systems to close the contactor and open the shielding gas solenoid valve (gas-shielded FCAW process only) when welding is started. Because flux-cored wires are easily deformed by excessive feed roll pressure, knurled feed rolls are generally used in the FCAW process. Some wire feeders use a single drive roll paired with an undriven pressure roll. Others have one or two

pairs of drive rolls. It is generally believed that systems having two pairs of drive rolls require the least drive roll pressure to provide dependable feeding. Both air-cooled and water-cooled welding guns are used in the semiautomatic FCAW process. Air-cooled guns are generally preferred, because they are simpler to maintain, lighter in weight, and less bulky. Water-cooled guns may be required when welding currents over 500 A are used, especially when the shielding gas is rich in argon. Air-cooled guns designed for gas-shielded welding should not be used for self-shielded welding, because the gun depends on the flow of shielding gas for proper cooling. Although curved-neck guns are the most common, straight guns are used to a limited extent. A trigger switch on the welding gun is closed to initiate wire feeding, welding current flow, and shielding gas flow. The electrode is delivered from the wire feeder to the gun through a flexible conduit. Standard conduit lengths are 3, 3.7, 4.6, and 6 m (10, 12, 15, and 20 ft). Other lengths may also be available. Mechanized and automatic FCAW equipment is not substantially different from that used in the semiautomatic

FCAW process. The power supply should be rated for 100% duty cycle. Power supplies capable of outputs up to 1000 A may be required for some applications. Constant-current systems are very seldom used for mechanized and automatic welding. The wire feed system is separated into a drive motor assembly and a welding control device, the latter of which often has a system to automatically start the travel mechanism when wire feed, current flow, and shielding gas flow are initiated. The welding control device is often equipped with a voltmeter and ammeter, as well. The welding gun in mechanized and automatic systems is often mounted directly to the drive motor assembly, eliminating the need for a conduit. The gun is usually straight, but curved-neck guns are also used. Both water- and air-cooled guns are used, depending on the welding current level and shielding gas. Most of the commonly available air-cooled guns can be used at levels up to 500 A with CO2 shielding gas. When argonrich shielding gas is used, the same gun may only be suitable for use at levels up to 300 A. Water-cooled guns are generally used at higher current levels. Various travel mechanisms are used, depending on the applications. These mechanisms include side-beam carriages, tractor-type carriages, and robots. Fume-Removal Equipment. In many cases, the amount of fumes generated by the FCAW process is sufficient to

require fume-removal equipment. Although such equipment can be as simple as exhaust fans in the shop roof, local fume collection is often necessary. These systems can be either collection hoods (located above the welding gun) or fumeextractor guns. These guns are more efficient at collecting fumes, but are heavier and more bulky than standard welding guns. Fume collection hoods must be repositioned each time the welding location is moved in order to be effective.


1. WELDING HANDBOOK, 8TH ED., VOL 2, AWS, 1991 Flux-Cored Arc Welding David W. Meyer, The Esab Group, Inc.

Base Metals Most weldable carbon and low-alloy steels can be welded by the FCAW process if suitable electrodes are available. Steels for which electrodes are available include: •

MILD STEELS, SUCH AS AISI 1010 TO 1030, ASTM A 36, A 285, A 515, AND A 516

• • • • • •

WEATHERING STRUCTURAL STEELS, SUCH AS ASTM A 588 HIGH-STRENGTH, LOW-ALLOY STEELS, SUCH AS ASTM A 710 HIGH-TEMPERATURE CHROMIUM-MOLYBDENUM STEELS, SUCH AS ASTM A 387 GRADES 12 (1CR-0.5MO) AND 22 (2.25CR-1MO) NICKEL-BASE STEELS, SUCH AS ASTM A 203 HIGH-STRENGTH QUENCHED AND TEMPERED STEELS, SUCH AS HY-80 AND ASTM A 514 AND A 517 MEDIUM-CARBON, HEAT-TREATABLE, LOW-ALLOY STEELS, SUCH AS AISI 4130

Several grades of stainless steel are welded with both gas-shielded and self-shielded flux-cored electrodes. Electrodes designed for welding AISI types 304, 316, 347, and others are available. Also available are electrodes suitable for joining stainless steels to carbon and low-alloy steels. Electrodes that can be used to weld some nickel-base alloys have also been introduced. Among the nickel-base alloys that have been joined using the FCAW process are alloy 600 (UNS N06600) and alloy 625 (UNS N06625). Some cast irons are also welded using nickel-base flux-cored electrodes designed specifically for this purpose. Flux-Cored Arc Welding David W. Meyer, The Esab Group, Inc.

Electrode Manufacture Flux-cored electrodes are generally manufactured using the process shown in Fig. 3. A flat sheath material is first formed into a "U" shape. The core ingredients are poured into this "U" at the desired rate. The sheath is then closed around the core materials to form a round tube. The diameter of this tube is then reduced, generally by drawing or rolling operations, to compress the core materials and bring the electrodes to a size that is usable for welding. The finished wire is than wound on spools, coils, or other packages. Other manufacturing methods, which are generally considered proprietary, are also used.

FIG. 3 FLUX-CORED ELECTRODE MANUFACTURING PROCESS. SOURCE: REF 2

The usability of a flux-cored electrode, as well as the properties and chemistry of the weld it deposits, primarily depends on the ingredients in the electrode core. Common core ingredients and their functions are listed in Table 1.

TABLE 1 FUNCTIONS OF COMMON CORE INGREDIENTS IN FCAW ELECTRODES CORE INGREDIENT

RUTILE (TIO2) FLUORSPAR (CAF2) LIME (CACO3) FELDSPAR SYNTHETIC FRITS MANGANESE SILICON TITANIUM ALUMINUM CHROMIUM, NICKEL, MOLYBDENUM

GAS FORMER

DEOXIDIZER

DENITRIFIER

X

X X X X

SLAG FORMER

VISCOSITY CONTROL

X X X X X

X X

X

ARC STABILIZER

ALLOY

X X X X X

X X X


2. C.E. JACKSON, "FLUXES AND SLAGS IN WELDING," WRC BULLETIN 190, WELDING RESEARCH COUNCIL Flux-Cored Arc Welding David W. Meyer, The Esab Group, Inc.

Electrode Diameters Flux-cored electrodes are produced in diameters ranging from 0.8 to 3.2 mm (0.030 to

1 in.). Electrodes for all-position 8

welding may be available in 0.8 mm (0.030 in.), 0.9 mm (0.035 in.), 1.2 mm (0.045 in.), 1.4 mm (0.052 in.), and 1.6 mm 1 1 5 in.) diameters. Electrodes for flat- and horizontal-position welding may be available in 1.6 mm ( in.), 2.0 mm ( 16 16 64 3 7 1 in.), 2.4 mm ( in.), 2.8 mm ( in.), and 3.2 mm ( in.) diameters. Other diameters may be made available by 32 64 8

(

agreement between the manufacturer and the purchaser. Electrode Classification Carbon and Low-Alloy Steel Electrodes. Carbon steel flux-cored electrodes are classified by AWS A5.20, "Specification for Carbon Steel Electrodes for Flux Cored Arc Welding." The layout of this classification system is defined in Fig. 4; an "X" indicates a position where a designator would be. For example, the designator indicating minimum tensile strength can be a 6, to denote 62 ksi (430 MPa), or a 7, to denote 72 ksi (495 MPa). The assignment of the final designator depends on the shielding requirements and features of the electrode, as well as the recommended welding polarity. These are described in Table 2. The designations EXXT-G and EXXT-GS are designed to allow the classification of electrodes that are not covered by any of the other classifications.

TABLE 2 USABILITY TYPE DESIGNATORS FOR FLUX-CORED ELECTRODES TYPE

SHIELDING(A)

SINGLE OR MULTIPASS(B)

T-1

GAS

MULTIPASS SPRAYLIKE

T-2

GAS

SINGLE

T-3

SELF

SINGLE

T-4

SELF

T-5

GAS

MULTIPASS GLOBULAR 27 J AT -30 °C DCEP(D) (20 FT · LBF AT -20 °F)

T-6

SELF

MULTIPASS SPRAYLIKE

T-7

SELF

MULTIPASS GLOBULAR NO REQUIREMENT

T-8

SELF

MULTIPASS GLOBULAR 27 J AT -30 °C DCEN (20 FT · LBF AT -20 °F)

T-10 SELF

TRANSFER

SPRAYLIKE

IMPACT TOUGHNESS REQUIREMENT

27 J AT -18 °C DCEP (20 FT · LBF AT 0 °F) NO DCEP REQUIREMENT

SPRAYNO LIKE REQUIREMENT MULTIPASS GLOBULAR NO REQUIREMENT

SINGLE

POLARITY(C)

DCEP DCEP

27 J AT -30 °C DCEP (20 FT · LBF AT -20 °F)

GLOBULAR NO REQUIREMENT T-11 SELF MULTIPASS SPRAYNO LIKE REQUIREMENT T-G NOT MULTIPASS NOT NO SPECIFIED SPECIFIED REQUIREMENT TNOT SINGLE NOT NO GS SPECIFIED SPECIFIED REQUIREMENT

DCEN

DCEN DCEN

SPECIAL CHARACTERISTICS

LOW SPATTER, FULL SLAG COVERAGE LOW SPATTER, FULL SLAG COVERAGE, HIGH DEOXIDIZERS HIGH SPEED HIGH DEPOSITION, LOW PENETRATION, CRACK RESISTANT IMPROVED TOUGHNESS, CRACK RESISTANT, THIN SLAG IMPROVED TOUGHNESS, DEEP PENETRATION CRACK RESISTANT, GOOD SLAG REMOVAL IMPROVED TOUGHNESS, CRACK RESISTANT HIGH SPEED GENERAL PURPOSE NOT SPECIFIED

NOT SPECIFIED NOT NOT SPECIFIED SPECIFIED

Source: AWS A5.20

(A) GAS-SHIELDED ELECTRODES ARE CLASSIFIED USING CO2 SHIELDING; HOWEVER, AR-CO2 MIXTURES ARE COMMONLY USED IN PRACTICE. (B) MULTIPASS ELECTRODES ARE SUITABLE FOR SINGLE-PASS OR MULTIPASS WELDING; SINGLE-PASS ELECTRODES ARE SUITABLE FOR SINGLE-PASS WELDING ONLY.

(C) DCEP, DIRECT-CURRENT ELECTRODE POSITIVE; DCEN, DIRECT-CURRENT ELECTRODE NEGATIVE. (D) T-5 ELECTRODES ARE CLASSIFIED USING DCEP; HOWEVER, THEY ARE SOMETIMES USED WITH DCEN IN PRACTICE.

FIG. 4 CLASSIFICATION SYSTEM FOR CARBON STEEL FLUX-CORED ELECTRODES. THE LETTER "X" AS USED IN THIS FIGURE AND IN ELECTRODE CLASSIFICATION DESIGNATIONS IN AWS SPECIFICATION A5.20-79 SUBSTITUTES FOR SPECIFIC DESIGNATIONS INDICATED BY THIS FIGURE. SOURCE: REF 3

It should be noted that the Canadian Standards Association specification W48.5-M1990, "Carbon Steel Electrodes for Flux- and Metal-Cored Arc Welding," includes two additional electrode types that are planned for inclusion in the next revision of the AWS A5.20 specification. These classifications are EXXT-9 and EXXT-12. T-9 electrodes are generally similar to T-1 electrodes, but with improved toughness. T-12 electrodes are generally similar to T-9 electrodes, but with a maximum tensile strength also specified. Low-alloy steel electrodes are classified under AWS A5.29, "Specification for Low Alloy Steel Electrodes for Flux Cored Arc Welding." This specification uses a classification system very similar to that used in specification AWS A5.20, except that a chemical composition designator is added to the designation. An electrode classified according to this specification will have the form EXXTX-X. All of the positions before the dash have the same meaning as those in specification AWS A5.20. The position behind the dash is the chemical composition designator, which consists of a letter and a number. The letter denotes the alloy type of the electrode as follows:

DESIGNATOR A B NI D W K G

ALLOY TYPE CARBON-MOLYBDENUM STEEL CHROMIUM-MOLYBDENUM STEEL NICKEL STEEL MANGANESE-MOLYBDENUM STEEL WEATHERING STEEL OTHER LOW-ALLOY STEELS NOT SPECIFIED

Specification AWS A5.29 classifies only the EXXT1-X, EXXT4-X, EXXT5-X, and EXXT8-X electrode types, and the usability designator has the name meaning as that in specification A5.20. Minimum tensile strengths of up to 830 MPa (120 ksi) are included in specification AWS A5.29. Impact toughness requirements are based on the strength, usability, and chemical composition requirements of the electrode.

Stainless steel electrodes are classified under AWS A5.22, "Specification for Flux Cored Corrosion Resisting

Chromium and Chromium-Nickel Steel Electrodes." Classifications to this specification have the form EXXXT-X. The first three positions are the chemical composition designator, which corresponds to the American Iron and Steel Institute (AISI) designations (such as 308, 316, and 410) of steels having a similar composition. The final position is the shieldingtype designator. T-1 types are designed for use with CO2 or Ar-CO2 shielding gases (classification requires the use of CO2). T-2 types are designed for use with Ar-2O2 shielding gas. T-3 types are self-shielded. A T-G type is included for electrodes that are not covered by the other shielding-type designators. Nickel-Base Electrodes. A specification for nickel-base flux-cored electrodes has not yet been published. The

manufacturers of these electrodes should be consulted for information on their usability, composition, and properties.


3. "SPECIFICATION FOR CARBON STEEL ELECTRODES FOR FLUX CORED ARC WELDING," AWS A5.20-79, AWS, 1979 Flux-Cored Arc Welding David W. Meyer, The Esab Group, Inc.

References

1. WELDING HANDBOOK, 8TH ED., VOL 2, AWS, 1991 2. C.E. JACKSON, "FLUXES AND SLAGS IN WELDING," WRC BULLETIN 190, WELDING RESEARCH COUNCIL 3. "SPECIFICATION FOR CARBON STEEL ELECTRODES FOR FLUX CORED ARC WELDING," AWS A5.20-79, AWS, 1979 Gas-Tungsten Arc Welding Grant Ken-Hicken, Sandia National Laboratory

Introduction GAS-TUNGSTEN ARC WELDING (GTAW), also known as HeliArc, tungsten inert gas (TIG), and tungsten arc welding, was developed in the late 1930s when a need to weld magnesium became apparent. Russell Meredith (Ref 1) developed a welding process using the inert gas helium and a tungsten electrode to fuse magnesium. This joining method replaced riveting as a method of building aircraft with aluminum and magnesium components. The HeliArc welding has continued to this day with many refinements and name changes, but with no change in the fundamentals demonstrated by Meredith (Ref 1). The melting temperature necessary to weld materials in the GTAW process is obtained by maintaining an arc between a tungsten alloy electrode and the workpiece (Fig. 1). Weld pool temperatures can approach 2500 °C (4530 °F). An inert gas sustains the arc and protects the molten metal from atmospheric contamination. The inert gas is normally argon, helium, or a mixture of helium and argon.

FIG. 1 SCHEMATIC SHOWING KEY COMPONENTS AND PARAMETERS OF THE GTAW PROCESS. SOURCE: REF 2

References

1. R. MEREDITH, U.S. PATENT 2,274,631 2. S. KOU, WELDING METALLURGY, JOHN WILEY & SONS, 1987, P 10 Gas-Tungsten Arc Welding Grant Ken-Hicken, Sandia National Laboratory

Applications Gas-tungsten arc welding is used extensively for welding stainless steel, aluminum, magnesium, copper, and reactive materials (for example, titanium and tantalum). The process can also be used to join carbon and alloy steels. In carbon steels, it is primarily used for root-pass welding with the application of consumable inserts or open-root techniques on pipe. The materials welded range from a few thousandths of an inch to several inches in thickness. Advantages and Limitations Advantages of GTAW include (Ref 3): • • • • • •

PRODUCES HIGH-QUALITY, LOW-DISTORTION WELDS FREE OF THE SPATTER ASSOCIATED WITH OTHER METHODS CAN BE USED WITH OR WITHOUT FILLER WIRE CAN BE USED WITH A RANGE OF POWER SUPPLIES WELDS ALMOST ALL METALS, INCLUDING DISSIMILAR ONES GIVES PRECISE CONTROL OF WELDING HEAT

The GTAW process is applicable when the highest weld quality is required. It can be used to weld almost all types of metals. The operator has excellent control of heat input, and vision is not limited by fumes or smoke from the process. Limitations of GTAW include (Ref 4):

• •

•

PRODUCES LOWER DEPOSITION RATES THAN CONSUMABLE ELECTRODE ARC WELDING PROCESSES REQUIRES SLIGHTLY MORE DEXTERITY AND WELDER COORDINATION THAN GAS METAL ARC WELDING (GMAW) OR SHIELDED METAL ARC WELDING (SMAW) FOR MANUAL WELDING LESS ECONOMICAL THAN CONSUMABLE ELECTRODE ARC WELDING FOR THICK SECTIONS GREATER THAN 9.5 MM (

•

3 8

IN.)

PROBLEMATIC IN DRAFTY ENVIRONMENTS BECAUSE OF DIFFICULTY IN SHIELDING THE WELD ZONE PROPERLY

Additional problems with the process may include: • • • • •

TUNGSTEN INCLUSIONS IF THE ELECTRODE IS ALLOWED TO CONTACT THE WELD POOL CONTAMINATION OF THE WELD METAL, IF PROPER SHIELDING OF THE FILLER METAL BY THE GAS STREAM IS NOT MAINTAINED LOW TOLERANCE FOR CONTAMINANTS ON FILLER OR BASE METALS CONTAMINATION OR POROSITY, CAUSED BY COOLANT LEAKAGE FROM WATERCOOLED TORCHES ARC BLOW OR ARC DEFLECTION, AS WITH OTHER PROCESSES

Power supplies for GTAW are usually the constant-current type with a drooping (negative) volt-ampere (V-A) curve.

Saturable reactors and thyristor-controlled units are the most common. Advances in the electronics industry have readily been accepted in the welding community, resulting in sophisticated, lightweight power supplies. Transistorized direct current (dc) power supplies are becoming common, and the newer rectifier-inverter supplies are very compact and versatile. The inverter power supply consists of three converters:

• • •

60 HZ PRIMARY ALTERNATING CURRENT (AC) IS RECTIFIED TO DC DIRECT CURRENT IS INVERTED TO HIGH-FREQUENCY AC ALTERNATING CURRENT IS RECTIFIED TO DC (REF 5)

The inverter supplies can be switched from constant current to constant voltage for GMAW, resulting in a very versatile piece of equipment. The inverter-controlled power supplies are more stable and have faster response times than conventional silicon-controlled rectifier (SCR) power supplies. Figure 2 compares the response of an inverter-controlled arc welding machine and a thyristor-controlled welding machine.

FIG. 2 STARTING CURRENT WAVEFORMS OF TWO POWER SOURCES TO SHOW RELATIVE RESPONSE TIMES OF EACH SOURCE. (A) THYRISTOR-CONTROLLED SOURCE. (B) INVERTER-CONTROLLED SOURCE. FASTER RESPONSE OF INVERTER-CONTROLLED ARC WELDING MACHINE (2 MS TO GO FROM 0 TO 100 A) INDICATES A MORE STABLE ARC.

Torch Construction. The welding torch holds the tungsten electrode that conducts the current to the arc, and it provides a means of shielding the arc and molten metal. The major components of a typical welding torch are shown in Fig. 3.

FIG. 3 SCHEMATIC SHOWING EXPLODED VIEW OF KEY COMPONENTS COMPRISING A GTAW MANUAL TORCH

Welding torches rated at less than 200 A are normally gas-cooled (that is, the shielding gas flows around the conductor cable, providing the necessary cooling). Water-cooled torches are used for continuous operation or at higher welding currents and are common for mechanized or automatic welding (see the section "GTAW Process Variations" in this article). The cooling water may be supplied to the torch from a recirculating tank that uses a radiator or chiller to cool the water. Electrodes. The nonconsumable electrodes used in GTAW are composed of tungsten or alloys of tungsten. The most

common electrode is a 2% ThO2-W alloy (EWTh-2). This material has excellent operating characteristics and good stability. Thoria is radioactive, so care must be taken when sharpening electrodes not to inhale metal dust. The grindings are considered hazardous waste in some states, and disposal may be subject to environmental regulations. Lanthaniated (EWLa-1) and yttriated tungsten electrodes have the best starting characteristics in that an arc can be started and maintained at a lower voltage. Ceriated tungsten (EWCe-2) is only slightly better than the thoriated tungsten with respect to arc starting and melt-off rate. Any of the aforementioned electrodes produce acceptable welds. The easy starting of the lanthaniated electrode is a result of the lower work function (Ref 6) which allows it to emit electrodes readily at a lower voltage. Pure tungsten is used primarily in ac welding and has the highest consumption rate. Alloys of zirconium are also used. Tungsten electrodes are classified on the basis of their chemical composition (Table 1). Requirements for tungsten

electrodes are given in the latest edition of ANSI/AWS A5.12 ("Specification for Tungsten and Tungsten Alloy Electrodes for Arc Welding and Cutting"). The shape of the electrode tip can affect the resulting weld shape. Electrodes with included angles from 60 to 120° are stable and give good weld penetration depth-to-width ratios (Ref 7). Electrodes with smaller included angles (5 to 30°) are used for grooved weld joints to eliminate arcing to the part side walls.

TABLE 1 CLASSIFICATION OF ALLOYING ELEMENTS IN SELECTED TUNGSTEN ALLOY ELECTRODES FOR GTAW APPLICATIONS AWS CLASSIFICATION

COLOR(A)

ALLOYING ELEMENT

ALLOYING OXIDE

EWP EWCE-2 EWLA-1 EWTH-1 EWTH-2 EWZR-1 EWG

GREEN ORANGE BLACK YELLOW RED BROWN GRAY

... CERIUM LANTHANUM THORIUM THORIUM ZIRCONIUM NOT SPECIFIED(B)

... CEO2 LA2O3 THO2 THO2 ZRO2 ...

ALLOYING OXIDE, WT%

2 1 1 2 0.25 ...

Source: Ref 4

(A) COLOR MAY BE APPLIED IN THE FORM OF BANDS, DOTS, AND SO ON, AT ANY POINT ON THE SURFACE OF THE ELECTRODE. (B) MANUFACTURER MUST IDENTIFY THE TYPE AND NOMINAL CONTENT OF THE RAREEARTH OXIDE ADDITION. Wire feed systems are made from a number of components and vary from simple to complex. The basic system consists of a means of gripping the wire sufficiently to pull it from the spool and push it through the guide tube to the point of welding. Electronic switches and controls are necessary for the electric drive motor. The wire will be fed into the leading edge for cold wire feeds and into the trailing edge for hot wire feeds. Cables, hoses, and gas regulators are necessary to deliver the process consumable of electricity, water, and inert gas

to the welding torch. Arc oscillation is used in both manual and mechanized welding (see the section "GTAW Process Variations" in this

article). The benefits in manual welding are basic to the control of the weld when adapting to changes in the weld joint and gap. In mechanized welding, the oscillation is typically produced by moving the entire welding torch mechanically or by moving the arc plasma with the aid of an externally applied magnetic field. Oscillation allows the welding heat to be placed at precise locations. This is advantageous when welding irregularly shaped parts. The number of welding passes and total heat input can be decreased when arc oscillation is used, because it reduces the cost as well as the weld shrinkage and upsetting. Figure 4 shows the effect of magnetic oscillation on distortion. Some alloys need the tempering produced by multipass stringer bead welding techniques, and arc oscillation should not be used. Externally applied magnetic fields can be used to stabilize the arc, minimize arc blow, and displace the arc plasma forward of the welding torch. This results in improved weld appearances and increased welding speeds in tube mills and other high-speed applications (Ref 9).

FIG. 4 EFFECT OF MAGNETIC OSCILLATION ON WELD DISTORTION IN 25 MM (

1 IN.) THICK STAINLESS 2

STEEL. (A) WELD PRODUCED WITH ARC OSCILLATION. (B) STRINGER BEADS PRODUCED IN WELD WITHOUT ARC OSCILLATION. NOTE MINIMAL UPSETTING WITH ARC OSCILLATION. SOURCE: REF 8


3. T. MYERS, WHY THE GROWING INTEREST IN GAS TUNGSTEN ARC WELDING? THE FABRICATOR, VOL 22 (NO. 9), NOV 1992, P 38 -39 4. WELDING HANDBOOK, 8TH ED., VOL 2, AMERICAN WELDING SOCIETY, 1991, P 74-107 5. T. BYRD, INVERTER POWER SOURCES: AN EFFICIENT ALTERNATIVE, WELD. J., VOL 72 (NO. 1), JAN 1993, P 37-40 6. A.A. SADEK, M. USHIO, AND F. MATSUDA, EFFECTS OF RARE EARTH METAL OXIDE ADDITIONS TO TUNGSTEN ELECTRODES, METALL. TRANS. A, VOL 21, 1990, P 3221-3234 7. J.R. KEY, ANODE/CATHODE GEOMETRY AND SHIELDING GAS INTERRELATIONSHIPS IN GTAW, WELD. J., VOL 59 (NO. 12), DEC 1980, P 364-S TO 370-S 8. G.K. HICKEN, N.D. STUCK, AND H.W. RANDAL, APPLICATIONS OF MAGNETICALLY CONTROLLED ARCS, WELD. J., VOL 55 (NO. 4), 1976, P 264-267 9. G.K. HICKEN, C.E. JACKSON, THE EFFECTS OF APPLIED OF MAGNETIC FIELDS ON WELDING ARCS, WELD. J., VOL 45 (NO. 11), NOV 1966, P 515-S TO 524-S Gas-Tungsten Arc Welding Grant Ken-Hicken, Sandia National Laboratory

Process Parameters Welding Current. Current is one of the most important operating conditions to control in any welding operation,

because it is related to the depth of penetration, welding speed, deposition rate, and quality of the weld. Fundamentally, there are but three choices of welding current:

• • •

DIRECT CURRENT ELECTRODE NEGATIVE (DCEN) DIRECT CURRENT ELECTRODE POSITIVE (DCEP) ALTERNATING CURRENT

Figures 5 and 6 show the effect of dc and ac on weld shape. Table 2 gives the recommended current relative to workpiece material.

TABLE 2 SUITABILITY OF TYPES OF CURRENT FOR GTAW OF SELECTED METALS

METAL WELDED

ALTERNATING DCEN DCEP CURRENT(A)

LOW CARBON STEEL: 0.38-0.76 MM (0.015-0.030 IN.)(A) G(B) 0.76-3.18 MM (0.030-0.125 IN.) NR

E E

NR NR

G(B) G(B) G(B) G(B) NR

E E E E E

NR NR NR NR NR

E E E G(B)

NR(C) NR(C) NR(C) E

G NR NR NR

G(B) NR NR

E E E

NR NR NR

E

NR(C)

G

E

NR(C)

NR

E G(B) NR

NR(C) E E

NR NR NR

HIGH-CARBON STEEL CAST IRON STAINLESS STEEL HEAT-RESISTANT ALLOYS REFRACTORY METALS ALUMINUM ALLOYS: 0.64 MM (0.025 IN.) >0.64 MM (0.025 IN.) CASTINGS BERYLLIUM COPPER AND ALLOYS: BRASS DEOXIDIZED COPPER SILICON BRONZE MAGNESIUM ALLOYS: ≤

3.2 MM (

>4.8 MM (

1 8

3 16

IN.) IN.)

CASTINGS SILVER TITANIUM ALLOYS

Note: E, excellent; G, good; NR, not recommended.

(A) STABILIZED. DO NOT USE ALTERNATING CURRENT ON TIGHTLY JIGGED ASSEMBLIES. (B) AMPERAGE SHOULD BE ABOUT 25% HIGHER THAN WHEN DCEN IS USED. (C) UNLESS WORK IS MECHANICALLY OR CHEMICALLY CLEANED IN THE AREAS TO BE WELDED

FIG. 5 EFFECT OF POLARITY ON GTAW WELD CONFIGURATION WHEN USING DIRECT CURRENT. (A) DCEN. DEEP PENETRATION, NARROW MELTED AREA, APPROXIMATE 30% HEAT IN ELECTRODE AND 70% HEAT IN BASE METAL. (B) DCEP. SHALLOW PENETRATION, WIDE MELTED AREA, APPROXIMATE 70% HEAT IN ELECTRODE AND 30% HEAT IN BASE METAL

FIG. 6 WELD CONFIGURATION AS A FUNCTION OF TYPE OF CURRENT (AC OR DC) USED

Alternating current is characterized as reversing the polarity of the work and electrode at 60 Hz. The rapidly changing polarity gives a cathodic cleaning action that is beneficial for oxide removal when welding aluminum and magnesium. The alternating currents result in electrode heating during the DCEP portion of each cycle. This necessitates the use of larger-diameter electrodes, normally made of pure tungsten. Variable polarity welding allows the frequency of polarity switching to be preset. This can produce the cleaning effects to ac welding and the high efficiency of dc welding. Direct current electrode negative is most often used in the GTAW process. This results in maximum application of heat to the work and maximum melting of the workpiece. Pulsed versus Nonpulsed Current. Nonpulsed or continuous current is the standard for GTAW. However, there are several advantages to using pulsed current. Pulsing produces the maximum amount of penetration while minimizing the total heat applied to the part. Pulsing also aids in timing the motion necessary in manual welding and allows the weld pool to cool between pulses. Microwelding refers to a class of weldments that are made at welding currents from 1 to 20 A. In most cases, the welding

is used for electronic applications, bellows, wires, and other components where heat input must be precisely controlled. Shielding Gases. The original GTAW process used helium as the shielding gas for welding magnesium and aluminum.

Today, argon is the predominant shielding gas. Argon is the least expensive of the inert gases used for shielding gas-tungsten arc welds, which is only partially

responsible for its widespread use. Argon has a low ionization potential (2.52 × 10-18 J, or 15.7 eV), making it easier to form an arc plasma than with other shielding gases. Argon is approximately 1.4 times heavier than air, so it displaces air, resulting in excellent shielding of the molten weld pool. Helium has an ionization potential of 3.92 × 10-18 J (24.5 eV), which results in more difficult arc initiation and operation

at a higher arc voltage. The higher arc voltage, V, results in a higher heat input, Q, for a given arc length and current, I:

Q = IVT

(EQ 1)

where Q is in Joules, I is in amperes; and t is in seconds. This high heat input can be very beneficial when welding copper, aluminum, and other high-conductivity materials. Helium shielding used with DCEN is very effective for welding thick aluminum. Gas Purity. Most materials can be welded using a welding grade torch gas with a purity of 99.995% or 50 ppm

impurities. However, some reactive materials (for example, titanium, molybdenum, and tantalum) require that the contaminant level be less than 50 ppm, which may require certified purity or the use of gas filters and purifiers (Ref 10). Gas Flow Rates. Helium, because of its low density, must be used at higher flow rates than argon. Typical flow rates for

argon are 7 L/min (15 ft3/h) and 14 L/min (30 ft3/h) for helium. Backup Purge. Protecting the molten weld pool from the atmosphere is very important in GTAW. Atmospheric

contamination can result in weld cracks, porosity, scaling, and an unacceptable granular appearance. The gas cup on the welding torch is the primary outlet of shielding gas for most GTAW applications. Back side shielding is important because the presence of oxygen can reduce weld metal penetration and result in the effects mentioned above (Ref 11). Balloon and water-soluble paper dams are sometimes used to minimize the volume to be purged. Copper backing bars and ceramics are sometimes used to hold shielding gas against the back surface of the molten weld and support the molten underbead. Reactive materials and special applications may require more elaborate shielding. This can be in the form of a simple trailing device providing the inert shielding gas or may be as elaborate as a special welding chamber equipped with gas purifiers and analyzers. Specially constructed plastic bags have been used successfully to weld large, irregularly shaped components (Fig. 7).

FIG. 7 PLASTIC BAG ENCLOSURE THAT SIMULATES A GLOVE BOX, USED TO PURGE IRREGULARLY SHAPED COMPONENTS IN GTAW OPERATIONS

Filler Metals. The thickness of the part to be welded will determine the need for filler metal additions. Material thinner than 3.2 mm (0.125 in.) can be successfully welded without filler metal additions. Filler metal, when needed, can be added manually in straight length or automatically from a roll or coil. The filler metal is normally added cold; hot wire can be used for automatic applications (Fig. 8). A welding insert is preplaced filler material of several possible configurations to aid in root-pass welding.

FIG. 8 SCHEMATIC SHOWING KEY COMPONENTS AND PARAMETERS OF A GTAW HOT WIRE SYSTEM. SOURCE: REF 4

Rods. Straight lengths of filler wire, typically 915 mm (36 in.) in length, are used for manual welding. Most straight

lengths are round in cross section, but some aluminum fillers are somewhat rectangular. Cold Wire. Coiled wire may be acquired in small 100 mm (4 in.) spools. Larger 305 mm (12 in.) spools or large coils can

weigh over 225 kg (500 lb). The larger coils are normally used with GMAW because it requires larger quantities of filler metal. The filler wire is fed into the leading edge of the weld pool during cold-wire welding. Hot wire GTAW utilizes a heated filler metal to increase the deposition rate of the process (Fig. 9). The wire is resistance heated to near the melting temperature and fed into the trailing edge of the weld pool. Deposition rates to 29 kg/h (65 lb/h) are achievable. The higher deposition rates obtained with hot wire make the process competitive for welds and overlays and improve productivity. Table 3 lists typical hot wire parameters.

TABLE 3 TYPICAL PARAMETERS FOR AUTOMATIC HOT WIRE GTAW ARC CURRENT, A

300 400 500

ARC VOLTAGE, V

TORCH TRAVEL SPEED

WIRE FEED RATE(A)

DEPOSITION RATE

10-12 11-13 12-15

mm/min 100-255 150-355 205-510

mm/min 2790-9400 4700-11300 7490-16900

kg/h 1.4-4.5 2.3-5.4 3.6-8.2

Note: Using a 4.0 to 4.8 mm (

(A)

in./min 4-10 6-14 8-20

in./min 110-370 185-445 292-665

5 3 to in.) diameter 2% Th tungsten electrode with a 75He-25Ar shielding gas. 32 16

WIRE DIAMETER: 1.14 MM (0.045 IN.)

lb/h 3-10 5-12 8-18

FIG. 9 DEPOSITION RATES FOR GTAW WITH COLD AND HOT FILLER WIRE ON A STEEL WORKPIECE

Welding inserts are used to produce a smooth uniform underbead. The insert is normally a separate piece of material, although integral inserts are sometimes used. Smooth underbead can be produced in a mechanized system or in manual systems when highly skilled welders are available. The insert can be obtained in several different configurations (Fig. 10). The purpose of the insert is to preplace the filler metal at the joint root. Inserts with compositions that differ from that of the base metal can be used to improve the weldability of some materials.

FIG. 10 SELECTED JOINT CONFIGURATIONS AND WELDING INSERTS USED IN GTAW PROCESSES. (A) Y-TYPE

5 IN., FOR INDEXING WHEN ACCESSIBILITY IS LIMITED; NO LAND REQUIRED. (B) COMPOUND 32 1 1 1 5 BEVEL WITH OPEN ROOT, IN. GAP; LAND, IN. THICK. WASHER-TYPE INSERT ( × IN.) SHOULD BE 8 16 8 32 1 USED IF INSERT REQUIRED. (C) INSIDE DIAMETER/OUTSIDE DIAMETER PREPARATION WITH OPEN ROOT, 8 3 1 5 IN. GAP; LAND, IN. THICK. WASHER-TYPE INSERT ( × IN.) SHOULD BE USED IF INSERT REQUIRED. 32 8 32 1 5 1 1 (D) WASHER-TYPE INSERT ( × IN.); LAND, IN. THICK. (E) J PREP MUSHROOM INSERT; LAND IN. 8 32 16 16 3 THICK; EXTENDED LAND, IN. LONG 32 INSERT,


4. WELDING HANDBOOK, 8TH ED., VOL 2, AMERICAN WELDING SOCIETY, 1991, P 74-107 10. K.F. KRYSIAK AND P.M. BHADHA, SHIELDING GAS PURIFICATION IMPROVES WELD QUALITY, WELD. J., NOV 1990, P 47-49 11. H. GIEPL, ORBITAL TIG (GTA) WELDING, STAINL. STEEL LEUR., MAY 1992, P 48-55 Gas-Tungsten Arc Welding Grant Ken-Hicken, Sandia National Laboratory

GTAW Process Variations Manual welding refers to the GTAW process in which the welder manipulates the welding torch by hand. If a motorized wire feeder is attached to the torch, the process is classified as semiautomatic welding.

Products generated by skilled manual welders account for a large proportion of GTAW applications. The equipment can be quite inexpensive, and properly trained welders can join a wide variety of materials. Manual welding is used extensively in stainless steel piping as well as for the root pass in carbon steel pipe welds. Mechanized welding may require adjustment to welding parameters in response to visual observation of the weld.

Machine or mechanized welding requires some specialized accessories. The basic system contains a means for holding and moving the welding torch as well as the workpiece. Because arc voltage is an essential variable in GTAW and is proportional to arc length, voltage feedback devices are often used with motorized torch holders to control the arc length. Narrow groove welding makes use of the GTAW cold wire welding process with a narrowed weld joint. Figure 11 shows a typical weld. Narrow groove welding is limited to mechanized welding applications where precise torch location can be maintained.

FIG. 11 TYPICAL NARROW GROOVE WELD PRODUCED IN MECHANIZED WELDING APPLICATIONS

Automatic welding does not require manual 0parameter adjustment or observation of the weld during the welding process. The most common application of automatic welding is associated with orbiting weld heads used to weld pipe and tubing. These devices attach to the workpieces and move around the circumference, fusing the metal. Most systems perform an autogenous weld; others have wire feed and oscillation capabilities. These systems are often used in conjunction with a computer to control the welding variables. Automatic controls utilizing microprocessors and computer numerical control servodrives make it possible to use one welding system to weld a variety of materials and shapes. Figure 12 shows 14 different parts welded with an automatic GTAW system.

WELDING SPECIFICATIONS FOR COMPONENTS SHOWN Welds were made with a 250 A water-cooled GTAW torch that used 1.6 mm (0.062 in.) diameter electrode 6061 ALUMINUM WORKPIECE(A) COMPONENT NO. 1

WORKPIECE TUBE DIAMETER, 22.22 (0.875) MM (IN.) WALL THICKNESS, 1.57 (0.062)

COMPONENT NO. 2

COMPONENT NO. 3

28.58 (1.125) 0.88 (0.035)

26.82 (1.056) 1.65 (0.065)

TYPE 304 STAINLESS STEEL WORKPIECE(B) COMPONENT COMPONENT NO. 4 NO. 5

12.7 (0.5)

19.1 (0.75)

0.63 (0.025)

0.63 (0.025)

MM (IN.) POWER SUPPLY PEAK CURRENT, A BACKGROUND, A PULSE RATE, PPS(C) WIDTH, % TORCH WELD TRAVEL SPEED, MM/MIN WELDING CYCLE TIME, S SHIELDING GAS FLOW RATE, L/MIN (FT3H) FILLER METAL DIAMETER, MM (IN.) FEED RATE, MM/MIN (IN./MIN) JOINT TYPE

50-100 15 5

52-60 15 6

70-85 15 6

20-45 15 8

35-45 15 8

65

65

65

75

75

405 (16)

405 (16)

405 (16)

380 (15)

380 (15)

8.0

4.5

6.5

6.0

10.0

8.5 (18)

8.5 (18)

8.5 (18)

7.1 (15)

7.1 (15)

0.8 (0.03)

0.8 (0.03)

0.8 (0.03)

...

...

3810 (150)

3560 (140)

3810 (140)

...

...

SADDLE

FILLET

FILLET/BUTT

FILLET/BUTT BUTT

(A) 4043 ALUMINUM FILLER METAL; AC/DC SQUARE WAVE POWER SUPPLY; 2% CERIATED ELECTRODE MATERIAL; 50% HE-50% AR SHIELDING GAS; AC PROCESS (B) NO FILLER METAL USED; DC PRECISION POWER SUPPLY; 2% THORIATED ELECTRODE MATERIAL; ARGON SHIELDING GAS; DC PROCESS (C) PULSES PER SECOND FIG. 12 TYPICAL COMPONENTS PRODUCED BY AUTOMATIC GTAW PROCESS. SOURCE: REF 12


12. K.J. PFAHL, AUTOMATIC CONTOUR WELDING OF TUBE AND PIPE, TUBE PIPE Q., VOL 3 (NO. 4), 1992, P 58-62 Gas-Tungsten Arc Welding Grant Ken-Hicken, Sandia National Laboratory

References

1. R. MEREDITH, U.S. PATENT 2,274,631 2. S. KOU, WELDING METALLURGY, JOHN WILEY & SONS, 1987, P 10 3. T. MYERS, WHY THE GROWING INTEREST IN GAS TUNGSTEN ARC WELDING? THE FABRICATOR, VOL 22 (NO. 9), NOV 1992, P 38 -39 4. WELDING HANDBOOK, 8TH ED., VOL 2, AMERICAN WELDING SOCIETY, 1991, P 74-107 5. T. BYRD, INVERTER POWER SOURCES: AN EFFICIENT ALTERNATIVE, WELD. J., VOL 72 (NO. 1), JAN 1993, P 37-40 6. A.A. SADEK, M. USHIO, AND F. MATSUDA, EFFECTS OF RARE EARTH METAL OXIDE ADDITIONS TO TUNGSTEN ELECTRODES, METALL. TRANS. A, VOL 21, 1990, P 3221-3234

7. J.R. KEY, ANODE/CATHODE GEOMETRY AND SHIELDING GAS INTERRELATIONSHIPS IN GTAW, WELD. J., VOL 59 (NO. 12), DEC 1980, P 364-S TO 370-S 8. G.K. HICKEN, N.D. STUCK, AND H.W. RANDAL, APPLICATIONS OF MAGNETICALLY CONTROLLED ARCS, WELD. J., VOL 55 (NO. 4), 1976, P 264-267 9. G.K. HICKEN, C.E. JACKSON, THE EFFECTS OF APPLIED OF MAGNETIC FIELDS ON WELDING ARCS, WELD. J., VOL 45 (NO. 11), NOV 1966, P 515-S TO 524-S 10. K.F. KRYSIAK AND P.M. BHADHA, SHIELDING GAS PURIFICATION IMPROVES WELD QUALITY, WELD. J., NOV 1990, P 47-49 11. H. GIEPL, ORBITAL TIG (GTA) WELDING, STAINL. STEEL LEUR., MAY 1992, P 48-55 12. K.J. PFAHL, AUTOMATIC CONTOUR WELDING OF TUBE AND PIPE, TUBE PIPE Q., VOL 3 (NO. 4), 1992, P 58-62 Plasma Arc Welding Ian D. Harris, Edison Welding Institute

Introduction PLASMA ARC WELDING (PAW) can be defined as a gas-shielded arc welding process where the coalescence of metals is achieved via the heat transferred by an arc that is created between a tungsten electrode and a workpiece. The arc is constricted by a copper alloy nozzle orifice to form a highly collimated arc column (Fig. 1). The plasma is formed through the ionization of a portion of the plasma (orifice) gas. The process can be operated with or without a filler wire addition.

FIG. 1 PLASMA ARC WELDING PROCESS, SHOWING CONSTRICTION OF THE ARC BY A COPPER NOZZLE AND A KEYHOLE THROUGH THE PLATE

Principles of Operation. Once the equipment is set up and the welding sequence is initiated, the plasma and shielding

gases are switched on. A pilot arc is then struck between a tungsten alloy electrode and the copper alloy nozzle within the torch (nontransferred arc mode), usually by applying a high-frequency open-circuit voltage. When the torch is brought in close proximity to the workpiece or when the selected welding current is initiated, the arc is transferred from the electrode to the workpiece through the orifice in the copper alloy nozzle (transferred arc mode), at which point a weld pool is formed (Fig. 1).

The PAW process can be used in two distinct operating modes, often described as the melt-in mode and the keyhole mode. The melt-in-mode refers to a weld pool similar to that which typically forms in the gas-tungsten arc welding (GTAW)

process, where a bowl-shaped portion of the workpiece material that is under the arc is melted. In the keyhole mode, the arc fully penetrates the workpiece material, forming a nominally concentric hole, or keyhole, through the thickness. The molten weld metal flows around the arc and resolidifies behind the keyhole as the torch traverses the workpiece. Current and Operating Modes. The PAW process uses three current modes: microplasma (melt-in mode), medium-

current plasma (melt-in mode), and keyhole plasma (keyhole mode). This categorization is primarily based on the level of welding current. The microplasma mode is usually defined in the current range from 0.1 to 15 A. The medium-current plasma mode ranges from 15 to 100 A. The keyhole plasma mode is above 100 A. There is a certain degree of overlap between these current ranges. For example, keyholing can be achieved at 70 A on a 2 mm (0.08 in.) sheet. Equipment is available for welding currents up to 500 A, although a 300 A maximum is typical. Microplasma and medium-current melt-in modes are used for material up to 3 mm (0.12, or

1 , in.) thick, whereas the keyhole plasma mode is used for 8

greater thicknesses and higher travel speeds. In addition to operating in a continuous and steady direct current electrode negative (DCEN) mode, the PAW process can be carried out using DCEN pulsed current, as well as in the variable polarity mode, which uses both direct current electrode positive (DCEP) and electrode negative polarity switching. The pulsed current mode (both DCEN and DCEN/DCEP) is most often used when current levels (typically, above 100 A) are employed for keyhole plasma welding. Pulsing the current widens the tolerance region of acceptance welding parameters, primarily by further stabilizing the formation of the keyhole itself (Fig. 2).

FIG. 2 TOLERANCE TO VARIATION IN WELDING CURRENT AND PLASMA GAS FLOW RATE IN PULSED AND CONTINUOUS CURRENT KEYHOLE WELDING; BOUNDARIES SHOW THE WELDING PARAMETER COMBINATIONS AT WHICH SPECIFIC DEFECTS ARE LIKELY TO OCCUR. WELDING PARAMETERS; NOZZLE BORE, 2.36 MM (0.0929 IN.); ELECTRODE DIAMETER, 4.8 MM (0.19 IN.)

The electrode positive component of the variable polarity plasma arc (VPPA) welding process promotes cathode etching of the tenacious surface oxide film when welding aluminum alloys, allowing good flow characteristics and consistent bead shape. Pulsing times are typically 20 ms for the electrode negative component and 3 ms for the electrode positive polarity. The VPPA welding process is used very effectively in specialized aerospace applications. The PAW process is generally applied when the high penetration of the keyhole welding mode can be exploited to minimize the number of welding passes and, hence, welding time. The time saved can reduce the direct labor element of the welding operation. At the other end of the scale, the microplasma operating mode is used to weld small, thin-section components (as low as 0.025 mm, or 1 mil, thick), where the high arc constriction and low welding current can be beneficial in controlling heat input and distortion. Advantages and Disadvantages. The advantages of the PAW process are primarily intrinsic to the keyhole mode of operation, because greater thicknesses of metal can be penetrated in a single pass, compared with other processes, such as GTAW. This greater amount of penetration allows a reduced amount of joint preparation. In some materials, for example, a square-grooved butt joint preparation can be used for thicknesses up to 12 mm (0.5 in.). The process can produce high weld integrity (similar to GTAW) while minimizing weld passes and, hence, welding times and labor costs. The columnar shape of the arc results in a greater tolerance to variations in torch standoff distance, when compared with the conical arc shape of a GTAW arc. The tungsten electrode used in the PAW process is protected from contamination by the constricting nozzle (Fig. 1). The longer arc length allows better viewing of the weld pool, which is important in manual welding.

Disadvantages include the greater capital equipment cost, when compared with its main rival, the GTAW process. Although high arc constriction achieves higher penetration, it also reduces the tolerance of the process to joint gaps and misalignment, when compared with the broader, conical arc of the GTAW process. The greater complexity of the PAW torch design and the greater number of parts requires more scheduled maintenance. The accurate set-back of the electrode tip, with respect to the nozzle orifice, is required to maintain consistent results. However, this task is facilitated by a general-purpose tool designed for nozzle removal and replacement and for electrode set-back adjustment. Plasma Arc Welding Ian D. Harris, Edison Welding Institute

Equipment A basic PAW system consists of a power source, a plasma control console, a water cooler, a welding torch, and a gas supply system for the plasma and shielding gases (Fig. 3).

FIG. 3 TYPICAL EQUIPMENT FOR PLASMA ARC WELDING

The power source, which supplies the main power for the welding system, is usually supplemented with a sequence controller and control console. The sequence controller sequences the timing of gas flow, arc initiation, main welding current control, and any up-slope and downslope parameters. In its simplest form, the plasma control console controls the gas flow for plasma and shielding gases from separate flow-meters and incorporates the high-frequency pilot arc initiation circuit. The welding torch can be manual or mechanized and is water cooled to avoid torch overheating and to maximize component life. In most PAW installations, plasma and shielding gases are supplied from separate gas cylinders, although bulk gas can readily be used. The gas supply is usually routed through the plasma control console, where the individual flow rates are set by the operator. The power source should be of a constant-current design. Transistorized power sources are most common, although

inverter power supplies are also available. It should have a minimum open-circuit voltage of 80 V to ensure the reliable initiation and transfer of the main arc current. The power source can be adjusted for welding current and it should have the capability to adjust the up slope and down slope of the current. It may be equipped with thumbwheels or potentiometers to select the parameters for pulse current operation, that is, peak and background current levels, as well as peak and background times. Welding Torches. Like those of the other arc welding processes, PAW torches are available in a range of sizes for

different power ratings and in manual and mechanized versions. The design principles are the same in each case. A tungsten alloy electrode is held in a collet within the torch body. To avoid one of the most common defects in plasma torches, it is critical to hold concentricity between the tungsten electrode and the orifice in the design and manufacture of the torch. The electrode assembly is set inside a plenum chamber and the plasma gas is supplied to this chamber. A

threaded copper alloy nozzle forms the front of this chamber and contains the nozzle orifice that is used to constrict the plasma arc. A shielding gas nozzle, usually of an insulating ceramic material, is threaded onto the front end of the torch and surrounds the constricting nozzle, creating an annulus through which the shielding gas is supplied. The torch is connected electrically to the power source and the electrode forms the negative pole of the circuit for dc welding. The gas hoses that supply the plasma and shielding gases and the water hoses that supply and remove water from the torch are all connected to the torch body or handle. These hoses are enclosed in a flexible sheath that extends from the torch to the components of the welding system. Most constricting nozzles have a single orifice in the center. However, multiple-nozzle orifices can be used with higherpower torches to achieve further arc constriction. The most common version of this type of nozzle has a central orifice flanked by a smaller orifice on each side. The common centerline of the three orifices is arranged at 90° to the weld line during operation. Electrodes. The nonconsummable electrode employed is usually a 2% thoriated tungsten electrode, that is, tungsten with

2% thorium oxide. The electrode specification is covered in AWS A5.12-92. The electrode size is selected according to the welding current level that will be used. The electrode is ground with a tapered point on the end, the angle of which depends on the selected welding current level. Electrode sizes and tip angles for a range of welding current levels are shown in Table 1.

TABLE 1 MAXIMUM CURRENT FOR PLASMA WELDING WITH SELECTED ELECTRODE DIAMETER, VERTEX ANGLE, AND NOZZLE BORE DIAMETER MAXIMUM CURRENT, A

ELECTRODE DIAM

mm MICROPLASMA TORCH RATING, 20 A 5 5 1.0 10 20 MEDIUM CURRENT TORCH RATING, 100 A 30 2.4 50 75 100 TORCH RATING, 200 A 50 4.8 100 160 200 TORCH RATING, 400 A 180 3.2 200

VERTEX ANGLE, DEGREES

in.

PLASMA(A) NOZZLE BORE DIAM

SHIELDING(B) SHROUD FLOW RATE DIAM

FLOW RATE

mm

in.

l/min gal/min mm in.

l/min gal/min

0.04

15

0.8 0.8 1.0

0.03 0.03 0.04

0.2 0.3 0.5

0.05 0.08 0.13

8

0.32 4-7

1-1.8

0.10

30

0.79 1.17 1.57 2.06

0.03 0.05 0.06 0.08

0.47 0.71 0.94 1.18

0.12 0.19 0.25 0.31

12

0.48 4-7

1-1.8

0.19

30

1.17 1.57 2.36 3.20

0.05 0.06 0.09 0.13

0.71 0.94 1.42 1.65

0.19 0.25 0.38 0.44

17

0.68 4-12

1-3.2

0.13

60(C)

2.82

0.11

2.4

0.63

18

0.72 2035

5.3-9.2

2.82(C) 0.11(C) 2.5

0.66

HIGH CURRENT TORCH RATING, 400 A 250 4.8

0.19

300 350 (A) (B) (C) (D)

60(C)

3.45(D) 0.14

3.0

0.79

3.45(D) 0.14(D) 3.5 3.96(D) 0.16(D) 4.1

0.92 1.08

...

...

2035

5.3-9.2

ARGON PLASMA GAS. ARGON AND AR-5H2 SHIELDING GAS. ELECTRODE TIP BLUNTED TO 1 MM (0.04 IN.) DIAMETER. MULTIPORT NOZZLE

The new types of tungsten electrodes, which contain oxides of rare-earth elements in place of the thorium oxide, can also be used. These electrodes have been shown to have greater tip life. However, they are more expensive and their usefulness in plasma arc welding may be limited because of the high level of protection provided by the nozzle. In lowcurrent microplasma welding applications, their better emissivity provides easier arc transfer and better overall performance. If a finer wire is required, to fit the joint or to avoid undercut at the weld toes, then wires suitable for gas-shielded arc welding (for example, those used for GTAW) should be employed. The appropriate AWS specification series is AWS A5.XX. Plasma (Orifice) and Shielding Gases. The plasma gas is used to generate the arc, whereas the shielding gas is used to provide the weld pool with supplementary shielding from atmospheric contamination while it solidifies and cools. The plasma gas is almost always argon. Gas properties affect both weld shape and quality. Flow rates, particularly of the plasma gas, are also important, because they control the extent of plasma constriction. The flow rate can vary from 0.1 L/min (0.026 gal/min) for microplasma welding up to 10 L/min (2.6 gal/min) for keyhole plasma welding. Considerable care is needed to regulate the gas flow rate if keyhole closure is required, because the flow rate must be sloped out to 1 to 2 L/min (0.26 to 0.52 gal/min) within about 1 s. Gas flow control is best achieved by electronic means.

The design and current rating of welding torches are based on argon plasma gas. Argon provides effective shielding, being heavier than air, and is cheaper than helium. Shielding gas selection is based on the type of base metal (Table 2).

TABLE 2 PLASMA AND SHIELDING GAS COMPOSITIONS

MATERIAL MILD STEEL

PLASMA GAS SHIELDING GAS ARGON ARGON ARGON-2-5% H2(A) LOW-ALLOY STEELS ARGON ARGON AUSTENITIC STAINLESS STEEL ARGON ARGON-2-5% H2 HELIUM(A) NICKEL AND NICKEL ALLOY ARGON ARGON ARGON-2-5% H2(A) TITANIUM ARGON ARGON 75HE-25AR(A) ALUMINUM AND ALUMINUM ALLOYS ARGON ARGON HELIUM(A) COPPER AND COPPER ALLOYS ARGON ARGON 75HE-25AR(A) (A)

ALSO USED

Helium and argon-helium mixtures can be used as shielding gases to increase the thermal conductivity of the gas and, hence, the heating effect on the weld pool. Helium results in wider weld pools than argon, because it produces a higher arc voltage. Hydrogen additions to argon shielding gas tends to promote slightly narrower weld pools through arc constriction and achieves a very clean weld pool appearance, because it is a reducing gas. Although helium and hydrogen can be added to argon in the shielding gas to give higher heat input, the use of gases with higher heat contents in the plasma gas can result in torch overheating and potential damage. Plasma Arc Welding Ian D. Harris, Edison Welding Institute

Applications Material Types. The PAW process is commonly used to weld stainless steels in a wide range of thicknesses. The process

can also be used with carbon and alloy steels, aluminum alloys, titanium alloys, copper and nickel alloys, and more specialized materials, such as zirconium and tantalum. The thicknesses that can be welded in a single pass range from 0.025 mm (1 mil) for microplasma applications to 12.5 mm (0.5 in.) for the VPPA welding of aluminum. Direct-current pulsing can be used on most materials. The PAW process is often carried out in an autogenous mode, that is, without filler wire. When edge beveling is used, a filler wire is required to complete the joint. A filler wire can also be used with the keyhole mode of operation to avoid undercut at high welding speeds. Wire composition depends on that of the parent materials in the joint. The same continuous-wound wire that is used in GTAW operations is suitable. The industries that use the PAW process can be categorized as those that weld thin-section sheet using microplasma or medium-current plasma welding and those that weld plate using keyhole plasma welding.

A wide range of small devices and assemblies made from thin stainless steel sheet, including bellows assemblies and associated fittings, are welded using the microplasma operating mode. The narrow weld bead that can be produced provides sheet-metal fabrications with a good cosmetic appearance. Furthermore, the high welding speed that can be achieved, coupled with the good tolerance to stand-off variations resulting from the columnar nature of the arc, makes the process attractive for high-volume production work. Microplasma, as well as medium-current plasma modes, can be used to spot weld guide wires and lamp filaments, as well as in other applications that require highly repetitive autogenous welds. This type of application allows a user to limit the number of high-frequency arc starts that would be required with tungsten-inert gas welding. Keyhole plasma welding is extensively used to weld stainless steel pipe and tankage. The process is applied to individual strakes from plate to make stainless steel vessels in the food and chemical processing industries. Circumferential welding of strakes also can be used to create these products. Longitudinal seam welds in stainless steel pipe with wall thickness of [ges]3 mm ([ges]0.12 in.) are ideally suited to keyhole plasma welding, because joint preparation is minimized and single-pass welding can be consistently achieved without the use of weld-backing devices. Pipes with wall thickness above 5 to 6 mm (0.20 to 0.24 in.) employ the keyhole mode of operation for the root pass. Depending on the material type and wall thickness, melt-mode PAW, GTAW, gas-metal arc welding (GMAW), or submerged arc welding (SAW) are used to complete the joint. High-alloy composition piping is similarly manufactured. The manufacturing of stainless steel tube from strip was one of the first applications of the PAW process. Because the process can reliably produce full-penetration welds without the use of backing, it is extensively used on tube mills, because a lack of access precludes welding from the inside. The VPPA welding process has been very effective for welding large space shuttle orbiter sections of aluminum alloys, particularly the external fuel tank. The cathodic cleaning action of the electrode positive portion of the current cycle breaks up tenacious surface oxide film on the aluminum. The process can be used on thicknesses up to 12.5 mm (0.5 in.) with square-grooved butt joints, and it can be used for the root pass on thicker sections with edge beveling. It is anticipated that this technique will be extensively employed in building aluminum alloys structures in a modular space station. Plasma Arc Welding

Ian D. Harris, Edison Welding Institute

Typical Components and Joints The most common joint configuration used with the PAW process is a butt joint. The microplasma mode is used with overlapped butt (micro-lap) joints and with joints that have integral weld metal as a result of flanged, butted edges on very thin metals (Table 3). Corner joints with edge welds are also commonly welded using the microplasma and mediumcurrent modes.

TABLE 3 JOINT CONFIGURATIONS FOR PLASMA WELDING SHEET AND TUBULAR COMPONENTS

Because the keyhole operating mode fully penetrates the workpiece, it is used exclusively on square-grooved butt joints and single-V joints with a root face (Table 3). For square-grooved butt joint preparation, the thickness that can be welded in a single pass depends on the fluid flow characteristics of the workpiece material for a given heat input from the plasma torch. Thus, alloys of titanium and zirconium can be butt welded with square-grooved preparation at greater thicknesses than steels and stainless steels. Generally, it is an industry-accepted practice to use square-grooved preparation without edge beveling for stainless steels up to 6 mm (0.24 in.) in thickness. A 4 or 5 mm (0.16 or 0.20 in.) root face and a single-V preparation, usually with a 60° included angle, are used on steels thicker than 6 mm (0.24 in.) (Table 4). The root is welded in the keyhole mode and the rest of the joint is completed with a filler wire addition by melt-mode PAW, GTAW, GMAW, or SAW, depending on the material type, joint volume to be filled, and the mechanical property requirements of the joint.

TABLE 4 KEYHOLE PLASMA WELDING CONDITIONS FOR PIPE PREFABRICATION IN TYPE 304L/316L STAINLESS STEEL

(a) Welded thickness, 5 mm (0.20 in.). (b) Welded thickness, 6 mm (0.24 in.)

In the PAW process, fixturing for part fit-up is more critical than it is for the GTAW process, primarily because of the narrower, more-constricted arc in PAW. The backing bar design and shielding gas techniques employed for GTAW are appropriate for microplasma and medium-current modes. During keyhole welding, however, the arc passes right through the workpiece, forming an efflux plasma (Fig. 1) that normally extends 10 mm (0.4 in.) below the back face of the joint. This characteristic must be accommodated when designing backing systems. Failure to allow sufficient gap can result in turbulence in the efflux plasma, causing weld pool disturbance and porosity. When welding longitudinal seams in tube or pipe, the weld can be started and stopped on run-on and run-off tabs, precluding the need for keyhole closure. However, when using the keyhole mode on circumferential seams of vessels or on butt joints in tubes or pipes, the keyhole will overlap the start of the weld to produce a complete joint. Keyhole closure without porosity has been an area of difficulty, particularly when butt welding tubes, and it represented an important concern for more-widespread use of the process. Modern electronic controls for simultaneous reduction in gas flow and current slope-out provide a more reliable method for keyhole closure. Plasma Arc Welding Ian D. Harris, Edison Welding Institute

Procedures Process Operating Procedure. Welding parameters, such as welding current, arc voltage, travel speed (for mechanized/automatic operation), and plasma and shielding gas flow rates, are set by the procedure and implemented by the welder. Torch parameters include the correct electrode vertex angle and set-back distance, as well as the correct orifice diameter for the welding current level.

The operating sequence for the PAW process was described in the section "Principles of Operation" at the beginning of this article. After the weld pool or keyhole is formed, the torch is traversed across the workpiece at the preset welding speed. Welding is terminated by the down slope of the welding current, with simultaneous sloping of the plasma gas flow rate if keyhole welding and keyhole closure are desired. Slope control for the keyhole mode is shown in Fig. 4.

FIG. 4 TYPICAL SLOPE CONTROL PATTERN FOR WELDING CURRENT AND PLASMA GAS FLOW WHEN STARTING AND CLOSING A KEYHOLE; EXAMPLE IS FOR 9.5 MM (

3 IN.) THICK STEEL 8

Most plasma arc welding is done in a mechanized or automated operation that does not require much operator intervention, except initially to set the parameters and to position the workpiece and torch. When current pulsing and keyhole closure operations are involved, numerous parameters must be set, which requires strict attention to detail, accurate part fit-up, and careful alignment of the torch relative to the joint. Inspection and Weld Quality Control. All of the common nondestructive evaluation (NDE) techniques are applicable

to plasma arc welds. Radiography and ultrasonic inspection are the most common techniques used. The fact that most plasma arc welding is carried out in the keyhole mode means that there is a penetration bead on the root side of the joint. Visual inspection for full penetration, as well as for correct width and profile of both the penetration bead and the weld surface profile, can readily be accomplished. Smooth, even underbeads can be achieved through good procedure development with tolerances improved through current pulsing. Troubleshooting. When troubleshooting a PAW operation, all of the previously defined basic parameters of the welding

schedule should be checked. Welding discontinuities can result from incorrect electrode set-back or from a worn or damaged nozzle orifice. The concentricity of the diameter of the nozzle orifice and the alignment of the electrode and nozzle orifice are very important. Worn nozzles should be replaced. The condition of the tungsten electrode tip should also be checked. An undercut is one of the most common defects in the PAW process. Depending on material type, thickness, and preset welding parameters, a two-sided undercut can occur above certain threshold welding speeds. This can only be eliminated by increasing welding current. A one-sided undercut can result from misalignment of the torch, the electrode/orifice and multiport nozzles, or from the mismatched fit-up of the workpieces. Correct alignment is essential. Plasma Arc Welding Ian D. Harris, Edison Welding Institute

Personnel Requirements The skill level and training requirements for the PAW process are generally similar to those for the GTAW process. When manual welding, the columnar nature of the plasma arc allows greater variation of the torch-to-workpiece distance without altering the size and shape of the weld pool. The fact that the electrode is protected from contamination by the nozzle means that tungsten contamination and tungsten inclusions in the weld metal are more readily avoided than they are in the GTAW process. From the viewpoint of an operator, disadvantages include the greater complexity of the torch,

in terms of nozzle orifice and electrode set-back considerations, and the correct setting of two gases (plasma and shielding gas). When mechanized welding is carried out, which is usually the case for the keyhole mode, the operator is required to set a greater number of parameters (compared to GTAW), particularly if current pulsing and keyhole closure operations are required. Specific operator training is required for pulsed keyhole operation, and particularly for VPPA operations, which are usually carried out in the vertical position, welding upward, and thus require a high level of operator competence. Computerization of welding parameter selection is the trend for VPPA welding and other mechanized or automated PAW operations. The health and safety issues related to the PAW process are very similar to those of other gas-shielded arc welding processes, especially GTAW. These include electrical shock; electromagnetic radiation hazards, particularly ultraviolet radiation; burns from hot metal parts; and welding fumes and gases, including ozone. The volume of welding fumes produced is low and is similar to that produced by the GTAW process. Hexavalent chromium and ozone are concerns when welding stainless steels and aluminum alloys, respectively. Like the GTAW process, the low level of general welding fumes associated with PAW results in a comparatively higher level of ozone. However, the rapid decay of ozone within a short distance of the arc, coupled with the high degree of mechanization and automation typical of the keyhole plasma welding operation, means that operator exposure is generally very low. The microplasma mode, which is more commonly used manually, employs such low current levels that fume and ozone levels are very low. Carbon Arc Welding Lance Soisson, Welding Consultants, Inc.

Introduction CARBON ARC WELDING (CAW) utilizes what is considered to be a nonconsumable electrode, made of carbon or graphite, to establish an arc between itself and either the workpiece or another carbon electrode. However, this electrode erodes fairly quickly and generates carbon monoxide (CO) gas that partially replaces the air around the arc, thereby providing the molten weld with some protection. The CAW process, which uses either single or twin electrodes, most closely resembles gas-tungsten arc welding (GTAW), where the arc is used only as a source of heat. The single-electrode arrangement usually operates with direct current (dc), electrode negative (straight polarity), using most dc power supplies. The twin-electrode arrangement usually operates with alternating current (ac), generally with small ac power supplies. Figure 1 shows typical configurations for both procedures.

FIG. 1 TYPICAL ARRANGEMENTS FOR SINGLE-ELECTRODE AND TWIN-ELECTRODE CARBON ARC WELDING

Although the CAW process has been almost completely replaced by the newer processes used in the welding industry, it is still used in certain applications. The process does produce adequate welds in thin-sheet steel, but should be used with caution in any critical application because it provides only limited shielding from the atmosphere. Carbon Arc Welding Lance Soisson, Welding Consultants, Inc.

Operation CAW electrodes can be either carbon or graphite, although baked carbon electrodes are most commonly used because they last longer and are more readily available. Electrode sizes range from 3.2 to 22 mm (

1 7 to in.). The typical torch is 8 8

air cooled for the smaller electrodes, but water cooled for the larger ones. The recommended current ranges for both electrode types are given in Table 1.

TABLE 1 RECOMMENDED CURRENT RANGES FOR CARBON AND GRAPHITE ELECTRODES

ELECTRODE DIAMETER CURRENT, A mm in. CARBON GRAPHITE SINGLE ELECTRODE 1 3.2 15-30 15-35 4.8 6.4 7.9 9.5 11.1 12.7 15.9 19 22.2

8 3 16 1 4 5 16 3 8 7 16 1 2 5 8 3 4 7 8

TWIN ELECTRODES 1 6.4 7.9 9.5

4 5 16 3 8

25-55

25-60

50-85

50-90

75-115

80-125

100-150

110-165

125-185

140-210

150-225

170-260

200-310

230-370

250-400

290-490

300-500

400-750

55

...

75

...

95-120

...

In the single-electrode operation, the carbon electrode is prepared to a long tapered point that is one half the

electrode diameter (see Fig. 1). The arc is established by "scratch starting," that is, by bringing the electrode into contact with the workpiece and immediately withdrawing it to the correct length for welding. The generally accepted arc length is maintained between 6.4 and 9.5 mm (

1 3 and in.). Holding the length too short can result in carburization of the base 4 8

metal, creating potentially brittle welds. This is particularly true when welding without filler metal. When the joint requires filler metal, the same technique that is used in GTAW or oxyfuel gas welding (OFW) should be employed. The welding rod should be fed into the weld pool with one hand while manipulating the arc with the other. If the arc is broken for any reason, then restart should not occur directly at the hot metal where welding stopped, because a hard spot could result. Rather, a restart should occur in front or to one side of the weld, followed by a quick return to the point where welding stopped. Care should be taken to either incorporate the restart area into the weld or to grind the arc strike after welding is completed. The arc strike areas can be very hard and can cause cracking if they are not fused into the weld or ground out. In the twin-electrode operation, the workpiece is not part of the electrical circuit. The heat for welding is produced

by creating an arc between the two electrodes. The electrode holder has two adjustable arms that grip the carbon electrodes at the proper length to maintain the arc. The operator maintains a constant distance between the carbon electrodes by adjusting them as they are consumed. An experienced welder can make these adjustments while still welding or brazing. Small ac machines are normally used in the twin-electrode process because the electrodes will wear at an equal rate. Direct current can be used as long as the positive-side electrode is one size larger than the negative electrode. The larger electrode compensates for the faster consumption rate on the positive side of the direct current. 1 3 to in.) and are copper coated for better current flow. The heat must be 4 8 1 maintained so that the copper does not melt more than 13 mm ( in.) back from the electrode tip. Only enough heat to 2

Electrodes range in size from 6.4 to 9.5 mm (

cause the filler metal to flow easily should be used (Table 1). High heat causes excessive electrode consumption and could cause carburization of the weld metal. Carbon Arc Welding Lance Soisson, Welding Consultants, Inc.

Applications The single-electrode carbon arc process is primarily used on steel, cast iron, and copper, although it can be used on most ferrous and nonferrous materials. With steels, it is principally applied to make outside corner welds on thinner-gage materials, where no filler metal is used. A good fit-up is required, and a fluxing agent is often used to promote better welds. The resulting welds arc smoother and, frequently, more economical than similar welds created by the shielded metal arc welding (SMAW) process. Another application is the welding of galvanized steel using a silicon bronze filler metal. In this instance, it is important to maintain a short arc and to direct the heat to the filler metal to minimize damage to the galvanic coating. Cast iron can be welded using the carbon arc process, but it is more often used in brazing. The casting must be preheated to approximately 650 °C (1200 °F) and then slowly cooled to reduce the possibility of cracking and to produce machinable welds. The process is also used on copper, because of the high heat capabilities of the larger electrodes. Direct current, electrode negative, should always be used when welding or brazing copper. The workpiece should be preheated to a temperature between 150 and 650 °C (300 and 1200 °F), or at least locally preheated in the weld area. Holding a long arc will allow the carbon to combine with oxygen to form the CO that somewhat shields the weld.

Twin carbon arc welding is used primarily as a maintenance tool in small shops, farms, or homes. It is most commonly applied in brazing or soldering operations, although it can be used to weld, preheat, or postheat smaller parts. Carbon Arc Welding Lance Soisson, Welding Consultants, Inc.

Selected References

• H.B. CARY, MODERN WELDING TECHNOLOGY, 2ND ED., PRENTICE-HALL, 1989 • WELDING HANDBOOK, 8TH ED., VOL 2, WELDING PROCESSES, AWS, 1991 Submerged Arc Welding Jonathan S. Ogborn, The Lincoln Electric Company

Introduction SUBMERGED ARC WELDING (SAW) is an arc welding process in which the arc is concealed by a blanket of granular and fusible flux (Ref 1, 2, 3). Heat for SAW is generated by an arc between a bare, solid-metal (or cored) consumablewire or strip electrode and the workpiece. The arc is maintained in a cavity of molten flux or slag, which refines the weld metal and protects it from atmospheric contamination. Alloy ingredients in the flux may be present to enhance the mechanical properties and crack resistance of the weld deposit.

References

1. R.L. O'BRIEN, WELDING HANDBOOK, VOL 11, AWS, 1991, P 191-232 2. D.L. OLSON ET AL., SUBMERGED ARC WELDING, VOL 6, 9TH ED., METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 114-152 3. THE PROCEDURE HANDBOOK OF ARC WELDING, THE LINCOLN ELECTRIC CO., 1973, P 6.3-1 TO 6.3-24 Submerged Arc Welding Jonathan S. Ogborn, The Lincoln Electric Company

Principles of Operation Figure 1 shows a typical setup for automatic SAW. A continuous electrode is being fed into the joint by mechanically powered drive rolls. A layer of granular flux, just deep enough to prevent flash through, is being deposited in front of the arc. Electrical current, which produces the arc, is supplied to the electrode through the contact tube. The current can be direct current (dc) with electrode positive (reverse polarity), with electrode negative (straight polarity), or alternating current (ac). Figure 2 shows the melting and solidification sequence of SAW. After welding is completed and the weld metal has solidified, the unfused flux and slag are removed. The unfused flux may be screened and reused. The solidified slag may be collected, crushed, resized, and blended back into new flux. Recrushed slag and blends of recrushed slag with unused (virgin) flux are chemically different from new flux. Blends of recrushed slag may be classified as a welding flux, but cannot be considered the same as the original virgin flux (Ref 4).

FIG. 1 SCHEMATIC SHOWING KEY COMPONENTS OF AUTOMATIC SUBMERGED ARC WELDING SETUP

FIG. 2 SCHEMATIC SHOWING TYPICAL WELD POOL DYNAMICS OF A SUBMERGED ARC WELD

Submerged arc welding is adaptable to both semiautomatic and fully automatic operation, although the latter, because of its inherent advantages, is more popular (see the section "Advantages of Submerged Arc Welding" in this article). In semiautomatic welding, the welder controls the travel speed, direction, and placement of the weld. A semiautomatic welding gun is designed to transport the flux and wire to the operator, who welds by dragging the gun along the weld joint. Semiautomatic electrode diameters are usually less than 2.4 mm (

3 in.) to provide sufficient flexibility and 32

feedability in the gun assembly. Manually guiding the gun over the joint requires skill because the joint is obscured from view by the flux layer. In automatic SAW, travel speed and direction are controlled mechanically. Flux may be

automatically deposited in front of the arc, while the unfused flux may be picked by a vacuum recovery system behind the arc. To increase deposition rate or welding speed, more than one wire can be fed simultaneously into the same weld pool. For example, in the twin arc process, two electrodes are fed into the same weld pool while sharing a common power source and contact tip. In tandem arc SAW, multiple electrodes are arranged with one in front of the other. Each electrode has an independent power supply and contact tip. The spacing, configuration, and electrical nature of the electrodes may be arranged to optimize welding speed and bead shape (Ref 5). Advantages of submerged arc welding include the following: •

• • • •

• • • •

THE ARC IS UNDER A BLANKET OF FLUX, WHICH VIRTUALLY ELIMINATES ARC FLASH, SPATTER, AND FUME (THUS MAKING THE PROCESS ATTRACTIVE FROM AN ENVIRONMENTAL STANDPOINT). HIGH CURRENT DENSITIES INCREASE PENETRATION AND DECREASE THE NEED FOR EDGE PREPARATION. HIGH DEPOSITION RATES AND WELDING SPEEDS ARE POSSIBLE. COST PER UNIT LENGTH OF JOINT IS RELATIVELY LOW. THE FLUX ACTS AS A SCAVENGER AND DEOXIDIZER TO REMOVE CONTAMINANTS SUCH AS OXYGEN, NITROGEN, AND SULFUR FROM THE MOLTEN WELD POOL. THIS HELPS TO PRODUCE SOUND WELDS WITH EXCELLENT MECHANICAL PROPERTIES. LOW-HYDROGEN WELD DEPOSITS CAN BE PRODUCED. THE SHIELDING PROVIDED BY THE FLUX IS SUBSTANTIAL AND IS NOT SENSITIVE TO WIND AS IN SHIELDED METAL ARC WELDING AND GAS METAL ARC WELDING. MINIMAL WELDER TRAINING IS REQUIRED (THUS, RELATIVELY UNSKILLED WELDERS CAN BE EMPLOYED). THE SLAG CAN BE COLLECTED, REGROUND, AND SIZED FOR MIXING BACK INTO NEW FLUX AS PRESCRIBED BY MANUFACTURERS AND QUALIFIED PROCEDURES.

Limitations of submerged arc welding include the following: • • • •

THE INITIAL COST OF WIRE FEEDER, POWER SUPPLY, CONTROLS, AND FLUXHANDLING EQUIPMENT IS HIGH. THE WELD JOINT NEEDS TO BE PLACED IN THE FLAT OR HORIZONTAL POSITION TO KEEP THE FLUX POSITIONED IN THE JOINT. THE SLAG MUST BE REMOVED BEFORE SUBSEQUENT PASSES CAN BE DEPOSITED. BECAUSE OF THE HIGH HEAT INPUT, SAW IS MOST COMMONLY USED TO JOIN STEELS MORE THAN 6.4 MM (

1 4

IN.) THICK.


4. "SPECIFICATION FOR CARBON STEEL ELECTRODES AND FLUXES FOR SUBMERGED ARC WELDING," A5.17-89, AWS, 1989, P 18 5. THE PROCEDURE HANDBOOK OF ARC WELDING, THE LINCOLN ELECTRIC CO., 1973, P 3.2-3, 6.3-21 TO 6.3-22, 6.3-49 TO 6.3-58 Submerged Arc Welding Jonathan S. Ogborn, The Lincoln Electric Company

Process Applications

If a steel is suitable for welding with gas-metal arc welding, flux-cored arc welding, shielded metal arc welding, or gastungsten arc welding, procedures can be developed to weld the steel with SAW (Ref 6). The main limitations of SAW are plate thickness and position. Because SAW is a high heat input and high deposition rate process, it is generally used to weld thicker-section steels. Although the welding of 1.6 mm ( over 6.4 mm (

1 in.) thick steel is possible, most SAW is done on plate 16

1 in.) in thickness. 4

Types of Metals. Submerged arc welding is most commonly used to join plain carbon steels. Alloy steels can be readily

welded with SAW if care is taken to limit the heat input as required to prevent damage to the heat-affected zone (HAZ). Low-heat-input procedures are available for welding alloy steels and heat-treated steels to prevent grain coarsening and cracking in the HAZ. Maintaining proper preheat and interpass temperature is also important when welding alloy steels to prevent weld metal and HAZ cracking and to develop the required mechanical properties in the weld deposit. Submerged arc welding can be used to join stainless steels and nonferrous alloys. It is also commonly used to produce a stainless or nonferrous overlay on top of a base metal. Joint Configurations. The most common weld deposits made with SAW are groove, fillet, and plug welds, and

surfacing deposits. For groove welds, the characteristic deep penetration of SAW plays a role in joint selection. Plate up 5 1 in.) thick can be completely welded from one side using a square butt joint with a 0.8 mm ( in.) root 8 32 5 opening. Beveled joints should be used for multiple-pass welds on plate thicknesses greater than 15.9 mm ( in.). 8

to 15.9 mm (

Backing bars should be used to prevent loss of flux, to prevent melt through, and to ensure full-penetration weldments on one-sided joints. Although groove welds are generally made in the flat position, it is possible to develop special procedures to weld horizontally. Single-pass fillet welds with throat sizes up to 7.9 mm (

5 in.) thick and multiple-pass fillet welds are usually made in 16

the horizontal position. Single-pass fillet welds over 7.9 mm (5/16 in.) thick are usually made in the flat position. Submerged arc plug welds are made in the flat position by puddling the electrode into the center of the hole until the weld is complete. Surfacing welds are made with both wire and strip electrode. Surfacing is done to rebuild worn parts with a wear-resistant material or to overlay a plain steel part with stainless steel or other alloys (Ref 7, 8, 9). To prevent porosity, the surface to be welded should be clean and free of grease, oil, paints, moisture, and oxides. All slag from tack welds or previous layers should be removed. Tack welds should be positioned so that the submerged arc weld completely melts out the tack. The workpiece should be firmly clamped to minimize distortion and the need for tack welds. Industrial Uses for SAW. Because SAW is used to join thick steel sections, it is primarily used for shipbuilding, pipe

fabrication, pressure vessels, and structural components for bridges and buildings. Other than joining, SAW is used to build up parts and overlay with stainless or wear-resistant steel (for example, rolls for continuous casting steel, pressure vessels, rail car wheels, and equipment for mining, mineral processing, construction, and agriculture).


6. THE PROCEDURE HANDBOOK OF ARC WELDING, THE LINCOLN ELECTRIC CO., 1973, P 6.3-26 TO 6.3-73 7. R.S. BROWN ET AL., ARC WELDING OF STAINLESS STEELS, VOL 6, 9TH ED., METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 327-329 8. K.C. ANTONY ET AL., HARDFACING, VOL 6, 9TH ED., METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 784 9. D.R. THOMAS ET AL., WELD OVERLAYS, VOL 6, 9TH ED., METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 804-819

Submerged Arc Welding Jonathan S. Ogborn, The Lincoln Electric Company

Equipment and Supplies The power supplied to the contact tip may be direct current or alternating current. In the United States, dc power is most common for welding applications that require less than 1000 A. In Europe and the Far East, however, ac is used extensively, even in applications requiring less than 1000 A. Currents over 1000 A dc tend to create arc blow problems. Alternating current is most commonly used for high-current applications, for applications where arc blow may be a problem, and in multiwire applications. Direct-current power supplies may be either transformer-rectifiers or motor/engine generators that produce either constant voltage or constant-current power. Alternating-current power sources are most commonly transformers. To produce a submerged arc weld, both flux and electrode are consumed. Each flux and electrode combination, along with the variation of base material and process parameters, will produce a unique weld deposit. Because the integrity of the weld deposit depends on these parameters, specific fluxes and electrodes must be used in combination to optimize the weld metal properties. Power Sources A constant-voltage power supply is self-regulating, so it can be used with a constant-speed wire feeder. No voltage or current sensing is required. The current is controlled by the wire diameter, the electrical stickout, and the wire-feed speed, while the voltage is controlled by the power supply. Constant-voltage dc power is the best choice for the high-speed welding of thin steel. Unlike constant voltage, constant-current power supplies are not self-regulating, so they must be used with voltagesensing variable-wire-feed speed controls. A constant-current wire feeder monitors arc voltage and adjusts the wire-feed speed in response to changes in the arc voltage. The wire-feed speed control attempts to maintain a constant arc length, while the power supply controls the arc current. The constant-current output of a conventional ac machine varies with time like a sine wave dropping through zero with each polarity reversal. The voltage associated with the current is approximately a square wave. Because the power output is zero with each polarity change, an open-circuit voltage greater than 80 V may be required to ensure arc initiation. The constant-current ac machine requires voltage-sensing variable-wire-feed speed controls. On newer, solid-state power supplies, the current and voltage output both approximate square waves, with the instantaneous polarity reversal reducing arc initiation problems. The solid-state power supplies have constant-voltage characteristics that may be used with constant-speed wire-feed controls. Flux Classification Fluxes can be categorized depending on the method of manufacture, the extent to which they can affect the alloy content of the weld deposit, and the effect on weld deposit properties. Classification Relative to Production Method. Based on the manufacturing process, there are two different types of

fluxes: fused and bonded. Fused Fluxes. The raw materials for a fused flux are dry mixed and melted in a furnace. The molten mixture is then

rapidly solidified, crushed, screened, and packaged. Because of their method of manufacture, fused fluxes typically do not contain ferroalloys and deoxidizers. Bonded Fluxes. The powdered ingredients of a bonded flux are dry blended, and then mixed with a binder, usually

potassium or sodium silicate. After pelletizing, the wet flux is dried in an oven or kiln, sized appropriately, and then packaged. The relatively low baking temperatures allow bonded fluxes to contain deoxidizers and ferroalloys. Classification Relative to Effect on Alloy Content of Weld Deposit. Independent of manufacturing method, a given

flux may be described as an active, neutral, or alloy flux, depending on its ability to change the alloy content of the weld deposit. With all submerged arc fluxes, variations in arc voltage and other welding variables will change the ratio of flux

consumed to electrode or weld metal deposited. This ratio is often referred to as the flux-to-wire ratio. Normal flux-towire ratios for SAW are 0.7 to 0.9. An increase in the flux-to-wire ratio may be caused by either an increase in arc voltage or a decrease in the welding current. Likewise, a decrease in the flux-to-wire ratio may be caused by a decrease in arc voltage or an increase in the welding current. How the weld deposit composition changes with voltage (flux-to-wire ratio) provides an additional means of describing a flux. Active fluxes contain controlled amounts of manganese and/or silicon. These alloys are added as ingredients in the flux to provide improved resistance to porosity and weld cracking caused by contaminants such as oxygen, nitrogen, and sulfur on the plate or in the plate composition itself. Active fluxes are primarily used to make single-pass welds. Because active fluxes contain deoxidizers such as manganese and silicon, the alloy in the weld metal will change with the flux-towire ratio. Changes in manganese and silicon content in the weld deposit will affect the strength and impact properties of the weld metal; therefore, the arc voltage must be more tightly controlled when welding with active fluxes than with neutral fluxes. Neutral fluxes contain little or no deoxidizers, and by definition, will not produce any significant change in the weld

metal composition as a result of a large change in arc voltage. Because neutral fluxes contain little or no alloy, they must rely on alloy in the electrode to provide deoxidation. Single-pass welds on oxidized plate may be prone to cracking or porosity due to insufficient alloy content and are usually best welded with active fluxes. Neutral fluxes are primarily used for welding multiple passes on plate exceeding 25 mm (1 in.) in thickness. To quantify the active/neutral behavior of a flux, the Wall neutrality number (N) was developed (Ref 10) and is defined as:

N = 100 [|∆%SI| + |∆%MN|]

(EQ 1)

where ∆%Si is the change in the weight percent of silicon, and ∆%Mn is the change in the weight percent of manganese. To determine the Wall neutrality number, two different weld deposit pads must be made. One weld pad is welded with the same parameters as those specified for the weld test plate for the flux and electrode being used. The other weld pad is welded with the same parameters except that the voltage is increased by 8 V. The ∆%Si and ∆%Mn quantities in Eq 1 are then determined by the difference in the amount of silicon and manganese in the two deposits. A flux electrode combination that produces a Wall neutrality number of 40 or less is considered neutral, while a number of 40 indicates an active flux. Alloy fluxes contain enough alloy as a flux ingredient to produce an alloy weld metal with a carbon steel electrode. They

are also used with alloy and stainless steel wire and strip electrodes. Alloy fluxes find application primarily in welding alloy steels and hardfacing. Because the alloy in the weld deposit is a function of arc voltage (flux-to-wire ratio), it is important to follow the manufacturer's recommended procedure for obtaining proper weld deposit composition. Classification Relative to Basicity Index. Aside from the manufacturing method and the active, neutral, or alloy

behavior of flux types, another common method used to describe submerged arc fluxes is the basicity index, BI. The basicity index is the ratio of strongly bound metallic oxides to weakly bound metallic oxides (Ref 11). The Boniszewski basicity index is defined by:

(EQ 2)

The basicity index is an estimate of the oxygen content in the weld metal and is therefore used to predict weld metal properties. Basic fluxes tend to have lower weld metal oxygen content with good weld metal toughness, while acidic fluxes tend to produce higher weld metal oxygen content and coarser microstructure with a lower resistance to cleavage. Fluxes with a basicity index greater than 1.5 are considered basic; fluxes below 1.0 are considered acidic. Acid fluxes are typically preferred for single-pass welding because of their superior operating and bead-wetting

characteristics. In addition, acid fluxes usually have more resistance than basic fluxes to porosity caused by plate contamination by oil, rust, and mill scale.

Basic fluxes tend to have better impact properties than acid fluxes. This advantage is particularly evident on large

multiple-pass welds. Highly basic fluxes generally produce weld metal with very good impact properties in large multiple-pass weldments. Basic fluxes tend to exhibit poorer welding characteristics than acid fluxes on single-pass welds. Their use should be limited to large multipass weldments where good weld metal notch toughness is required. Selection of Electrodes Relative to Steel and Alloy Type There are three different types of consumable electrodes: solid wire, cored wire, and strip. Electrodes for SAW are available for welding carbon steel, low-alloy steel, stainless steel, and nickel-base alloys. Wire diameters vary from 1.6 to 6.4 mm (

1 1 to in.). Carbon steel wire usually has a light copper coating, which protects the wire from corrosion and 16 4

provides good electrical contact in the welding tip. Cored electrodes were developed to provide a low-cost method of increasing the number of alloys that can be welded with SAW. Cored electrodes typically consist of a mild steel tube with alloying elements in the center. Compositions of cored electrodes vary from mild and low-alloy steels through tool steels to stainless steels and high-alloy austenitic steels. Strip electrodes are generally used for welding overlays and for hardfacing applications. The advantages of strip are wide weld beads, low penetration, low dilution, and high deposition rates. The equipment is very similar to SAW with wire except that special drive rolls are needed to feed the strip, and sometimes auxiliary magnetic fields are set up to improve bead shape and tie-in. Strip thicknesses vary from 0.5 to 1.0 mm (0.02 to 0.04 in.), and widths vary from about 25 to 100 mm (1 to 4 in.). Plain Carbon Steels. The requirements for the classification of electrodes and fluxes for welding carbon steels are

detailed in Ref 12. Compositional requirements for solid electrodes are given in Table 1. Requirements for composite electrodes are defined in Table 2. Classification of solid electrodes is based on the chemical composition of the electrode. Classification of composite electrodes is based on weld deposit composition. Flux classification is based on the soundness and mechanical properties of the weld metal produced with a particular flux and electrode combination. The classification system used to describe flux-electrode combinations for mild steel welding is defined in Fig. 3. Interpretation of the classification system code is illustrated in the flux-electrode combination examples that follow.

TABLE 1 COMPOSITION OF SOLID-CARBON STEEL ELECTRODES ELECTRODE AWS CLASSIFICATION

UNS NO.

COMPOSITION, WT%(A)(B) C Mn Si

LOW-MANGANESE ELECTRODES EL8 K01008 0.10 0.25/0.60 EL8K K01009 0.10 0.25/0.60 EL12 K01012 0.04/0.14 0.25/0.60 MEDIUM-MANGANESE ELECTRODES EM12 K01112 0.06/0.15 0.80/1.25 EM12K K11113 0.05/0.15 0.80/1.25 EM13K K01313 0.06/0.16 0.90/1.40 EM14K K01314 0.06/0.19 0.90/1.40 EM15K K01515 0.10/0.20 0.80/1.25 HIGH-MANGANESE ELECTRODES EH11K K11140 0.07/0.15 1.40/1.85 EH12K K01213 0.06/0.15 1.50/2.00 EH14 K11585 0.10/0.20 1.70/2.20

S

P

Cu(c)

Ti

0.07 0.030 0.030 0.35 . . . 0.10/0.25 0.030 0.030 0.35 . . . 0.10 0.030 0.030 0.35 . . . 0.10 0.10/0.35 0.35/0.75 0.35/0.75 0.10/0.35

0.030 0.030 0.030 0.025 0.030

0.030 0.030 0.030 0.025 0.030

0.35 0.35 0.35 0.35 0.35

... ... ... 0.03-0.17 ...

0.80/1.15 0.030 0.030 0.35 . . . 0.25/0.65 0.025 0.025 0.35 . . . 0.10 0.030 0.030 0.35 . . .

Source: Ref 12

(A) THE FILLER METAL SHALL BE ANALYZED FOR THE SPECIFIC ELEMENTS FOR WHICH VALUES ARE SHOWN IN THIS TABLE. IF THE PRESENCE OF OTHER ELEMENTS IS INDICATED, IN THE COURSE OF THIS WORK, THE AMOUNT OF THOSE ELEMENTS SHALL

BE DETERMINED TO ENSURE THAT THEIR TOTAL (EXCLUDING IRON) DOES NOT EXCEED 0.50%. (B) SINGLE VALUES ARE MAXIMUM. (C) COPPER LIMIT INCLUDES ANY COPPER COATING THAT MAY BE APPLIED TO THE ELECTRODE. TABLE 2 COMPOSITION OF CARBON STEEL WELD METAL MADE WITH COMPOSITE ELECTRODE

ELECTRODE AWS COMPOSITION, WT%(A)(B)(C) CLASSIFICATION C Mn Si S P Cu EC1 0.15 1.80 0.90 0.035 0.035 0.35 Source: Ref 12

(A) WELD METAL SHALL BE ANALYZED FOR THE SPECIFIC ELEMENTS FOR WHICH VALUES ARE SHOWN IN THIS TABLE. IF THE PRESENCE OF OTHER ELEMENTS IS INDICATED, IN THE COURSE OF THIS WORK, THE AMOUNT OF THOSE ELEMENTS SHALL BE DETERMINED TO ENSURE THAT THEIR TOTAL (EXCLUDING IRON) DOES NOT EXCEED 0.50%. (B) SINGLE VALUES ARE MAXIMUM. (C) A LOW-DILUTION AREA OF A GROOVE WELD OR A FRACTURED TENSION TEST SPECIMEN AS DESCRIBED IN AWS A5.17 (SEE REF 10) MAY BE SUBSTITUTED FOR THE WELD PAD, AND SHALL MEET THE ABOVE REQUIREMENTS. IN CASE OF DISPUTE, THE WELD PAD SHALL BE THE REFEREE METHOD.

FIG. 3 CLASSIFICATION SYSTEM APPLICATIONS. SOURCE: REF 12

FOR

CARBON

STEEL

ELECTRODES

AND

FLUXES

USED

IN

SAW

F7A6-EM12K is a complete designation. It refers to a flux that will produce weld metal that, in the as-welded condition, will have a tensile strength no lower than 485 MPa (70 ksi) and Charpy V-notch impact strength of at least 27 J (20 ft · lbf) at -51 °C (-60 °F) when produced with an EM12K electrode under the conditions called for in this specification. F7A4-EC1 is a complete designation for a flux when the trade name of the electrode used in classification is indicated as

well. It refers to a flux that will produce weld metal with that electrode, which in the as-welded condition, will have a tensile strength no lower than 485 MPa (70 ksi) and Charpy V-notch energy of at least 27 J (20 ft · lbf) at -40 °C (-40 °F) under the conditions called for in this specification.

Low-Alloy Steels. The requirements for low-alloy steel welding electrodes are detailed in Ref 13. The classification

system for flux-electrode combinations of low-alloy steel is shown in Fig. 4. Table 3 shows the AWS electrode classification as determined by the composition of the solid electrode composition, while Table 4 is the flux-electrode classification requirement for the weld deposit composition of both solid and composite flux-electrode combinations. Interpretation of the classification system code is illustrated in the flux-electrode combination examples that follow.

TABLE 3 COMPOSITION OF LOW-ALLOY STEEL SOLID ELECTRODES ELECTRODE AWS CLASSIFICATI ON

UNS NO.

EL12(E)

K0101 2

EM12K(E)

K0111 3

EA1

K1122 2

EA2

K1122 3

EA3

K1142 3

EA3K

K2145 1

EA4

K1142 4

EB1

K1104 3

EB2

K1117 2

EB2H

K2301 6

EB3

K3111 5

EB5

K1218 7

EB6(F)

S5028 0

EB6H

S5018 0

COMPOSITION, WT%(A)(B)(C) C Mn Si S

0.04 0.14 0.05 0.15 0.07 0.17 0.07 0.17 0.07 0.17 0.07 0.12 0.07 0.17 0.10

0.07 0.15 0.28 0.33 0.05 0.15 0.18 0.23 0.10

0.25 0.40

0.25 0.60 0.80 1.25 0.65 1.00 0.95 1.35 1.65 2.20 1.60 2.10 1.20 1.70 0.40 0.80 0.45 1.00 0.45 0.65 0.40 0.80 0.40 0.70 0.35 0.70 0.75 1.00

P

Cr

Ni

Mo

Cu(d)

V

Al

Ti

Zr

0.10

0.03 0

0.03 0

...

...

...

0.35

...

...

...

...

0.10 0.35 0.20

0.03 0

0.03 0

...

...

...

0.35

...

...

...

...

0.03 0

0.02 5

...

...

0.35

...

...

...

...

0.20

0.03 0

0.02 5

...

...

0.35

...

...

...

...

0.20

0.03 0

0.02 5

...

...

0.35

...

...

...

...

0.50 0.80 0.20

0.02 5

0.02 5

...

...

0.35

...

...

...

...

0.03 0

0.02 5

...

...

0.35

...

...

...

...

0.05 0.30 0.05 0.30 0.55 0.75 0.05 0.30 0.40 0.60 0.05 0.50 0.25 0.50

0.02 5

0.02 5

...

0.35

...

...

...

...

0.03 0

0.02 5

0.35

...

...

...

...

0.01 5

0.01 5

0.30

...

...

0.02 5

0.35

0.20 0.30 ...

...

0.02 5

...

...

...

0.02 5

0.02 5

0.30

...

...

...

...

0.02 5

0.02 5

0.35

...

...

...

...

0.03 0

0.02 5

0.40 0.75 1.00 1.75 1.00 1.50 2.25 3.00 0.45 0.65 4.50 6.50 4.80 6.00

0.45 0.65 0.45 0.65 0.45 0.65 0.40 0.60 0.45 0.65 0.45 0.65 0.45 0.65 0.40 0.65 0.90 1.10 0.90 1.20 0.45 0.70 0.45 0.65

0.35

...

...

...

...

...

...

...

...

...

...

ELECTRODE AWS CLASSIFICATI ON

UNS NO.

COMPOSITION, WT%(A)(B)(C) C Mn Si S

EB8(F)

S5048 0

0.10

ENI1

K1104 0

0.12

ENI2

K2101 0

ENI3

K3131 0

ENI4

K1148 5

ENIIK

K1105 8

EF1

K1116 0

EF2

K2145 0

EF3

K2148 5

EF4

K1204 8

EF5

K4137 0

EF6

K2113 5

EM2(G)

K1088 2

EM3(G)

K2101 5

EM4(G)

K2103 0

EW

K1124 5

EG

...

0.30 0.65

P

Cr

Ni

Mo

Cu(d)

V

Al

Ti

Zr

8.00 10.5 0 0.15

...

0.80 1.20

0.35

...

...

...

...

0.75 1.25 2.10 2.90 3.10 3.80 1.60 2.10 0.75 1.25 0.95 1.60 0.40 0.80 0.70 1.10 0.40 0.80 2.30 2.80 1.75 2.25 1.40 2.10 1.90 2.60 2.00 2.80 0.40 0.80

0.30

0.35

...

...

...

...

...

0.35

...

...

...

...

...

0.35

...

...

...

...

0.10 0.30 ...

0.35

...

...

...

...

0.35

...

...

...

...

0.25 0.55 0.40 0.65 0.40 0.65 0.15 0.30 0.45 0.65 0.40 0.65 0.25 0.55 0.25 0.65 0.30 0.65 ...

0.35

...

...

...

...

0.35

...

...

...

...

0.35

...

...

...

...

0.35

...

...

...

...

0.50

...

...

...

...

0.35

...

...

...

...

0.25

0.05

0.1 0

0.1 0

0.1 0

0.25

0.04

0.1 0

0.1 0

0.1 0

0.25

0.03

0.1 0

0.1 0

0.1 0

0.30 0.80

...

...

...

...

0.05 0.50

0.03 0

0.04 0

0.75 0.05 1.25 0.30 0.12 0.75 0.05 1.25 0.30 0.13 0.60 0.05 1.20 0.30 0.12 0.60 0.10 0.19 1.00 0.30 0.12 0.80 0.40 1.40 0.80 0.07 0.90 0.15 0.15 1.70 0.35 0.10 1.70 0.20 0.18 2.40 0.10 1.70 0.30 0.18 2.40 0.16 0.60 0.15 0.23 0.90 0.35 0.10 1.70 0.20 0.17 2.20 0.07 1.45 0.10 0.15 1.90 0.30 0.10 1.25 0.20 1.80 0.60 0.10 1.40 0.20 1.80 0.60 0.10 1.40 0.20 1.80 0.60 0.12 0.35 0.20 0.65 0.35 NOT SPECIFIED

0.02 0

0.02 0

0.02 0

0.02 0

...

0.02 0

0.02 0

0.15

0.02 0

0.01 5

...

0.02 0

0.02 0

...

0.02 5

0.02 5

...

0.02 5

0.02 5

...

0.02 5

0.02 5

...

0.03 5

0.02 5

0.40 -.60

0.01 0

0.01 0

0.01 5

0.01 5

0.01 0

0.01 0

0.25 0.50 0.20 0.55 0.30

0.01 0

0.01 0

0.55

0.01 0

0.01 0

0.60

0.04 0

0.03 0

0.50 0.80

Source: Ref 13

(A) THE FILLER METAL SHALL BE ANALYZED FOR THE SPECIFIC ELEMENTS FOR WHICH VALUES ARE SHOWN IN THIS TABLE. IF THE PRESENCE OF OTHER ELEMENTS IS INDICATED, IN THE COURSE OF THIS WORK, THE AMOUNT OF THOSE ELEMENTS SHALL BE DETERMINED TO ENSURE THAT THEIR TOTAL (EXCLUDING IRON) DOES NOT EXCEED 0.50 WT%. (B) SINGLE VALUES ARE MAXIMUM. (C) THE LETTER "N" AS A SUFFIX TO A CLASSIFICATION INDICATES THAT THE ELECTRODE IS INTENDED FOR WELDS IN THE CORE BELT REGION OF NUCLEAR REACTOR VESSELS. THIS SUFFIX CHANGES THE LIMITS ON THE PHOSPHORUS, VANADIUM, AND COPPER AS FOLLOWS: P = 0.012% MAXIMUM: V = 0.05% MAXIMUM; CU = 0.08% MAXIMUM. "N" ELECTRODES SHALL NOT BE COATED WITH COPPER OR ANY MATERIAL CONTAINING COPPER. THE "EF5" AND "EW" ELECTRODES SHALL NOT BE DESIGNATED AS "N" ELECTRODES. (D) THE COPPER LIMIT INCLUDES ANY COPPER COATING THAT MAY BE APPLIED TO THE ELECTRODE. (E) THE EL12 AND EM12K CLASSIFICATIONS ARE IDENTICAL TO THOSE SAME CLASSIFICATIONS IN ANSI/AWS A5.17-89 (REF. 10). THEY ARE INCLUDED IN THIS SPECIFICATION BECAUSE THEY ARE SOMETIMES USED WITH AN ALLOY FLUX TO DEPOSIT SOME OF THE WELD METALS CLASSIFIED IN TABLE 4. (F) THE EB6 AND EB8 CLASSIFICATIONS ARE SIMILAR, BUT NOT IDENTICAL, TO THE ER502 AND ER505 CLASSIFICATIONS, RESPECTIVELY, IN ANSI/AWS A5.9-80, "SPECIFICATION FOR CORROSION-RESISTING CHROMIUM AND CHROMIUM-NICKEL STEEL BARE AND COMPOSITE METAL CORED AND STRANDED ARC WELDING ELECTRODES AND WELDING RODS." THESE CLASSIFICATIONS WILL BE DROPPED FROM THE NEXT REVISION OF A5.9 (SEE REF 14). (G) THE COMPOSITION RANGES OF CLASSIFICATIONS WITH THE "EM" PREFIX ARE INTENDED TO CONFORM TO THE RANGES FOR SIMILAR ELECTRODES IN THE MILITARY SPECIFICATIONS. TABLE 4 COMPOSITION OF LOW-ALLOY STEEL WELD METAL (BOTH SOLID FLUX-ELECTRODE AND COMPOSITE FLUX-ELECTRODE COMBINATIONS) AWS WELD METAL CLASSIFICATI ON(E)

UNS NO.

COMPOSITION, WT%(A)(B)(C)(D)

SOLID

COMPOSI TE

C

Mn

Si

S

A1

K112 22

W17041

0.1 2

1.0 0

0.8 0

A2

K112 23

W17042

0.1 2

1.4 0

A3

K114 23

W17043

0.1 5

A4

K114 24

W17044

0.1 5

P

Cr

Ni

Mo

Cu

V

Ti

Z r

0.04 0.03 . . . 0 0

...

0.3 5

...

...

. . .

0.8 0

0.04 0.03 . . . 0 0

...

0.3 5

...

...

. . .

2.1 0

0.8 0

0.04 0.03 . . . 0 0

...

0.3 5

...

...

. . .

1.6 0

0.8 0

0.04 0.03 . . . 0 0

...

0.4 00.6 5 0.4 00.6 5 0.4 00.6 5 0.4 00.6 5

0.3 5

...

...

. . .

B1

K110 43

W51040

0.1 2

1.6 0

0.8 0

0.04 0.03 0.40 . . . 0 0 0.65

B2

K111 72

W52040

0.1 5

1.2 0

0.8 0

0.04 0.03 1.00 . . . 0 0 1.50

B2H

K230 16

W52240

1.2 0

0.8 0

0.04 0.03 1.00 . . . 0 0 1.50

B3

K311 15

W53040

0.1 00.2 5 0.1 5

1.2 0

0.8 0

0.04 0.03 2.00 . . . 0 0 2.50

B4

...

W53346

0.1 2

1.2 0

0.8 0

0.04 0.03 1.75 . . . 0 0 2.25

B5

K121 87

W51348

0.1 8

1.2 0

0.8 0

0.04 0.03 0.40 . . . 0 0 0.65

B6

S5028 W50240 0

0.1 2

1.2 0

0.8 0

0.04 0.03 4.50 . . . 0 0 6.00

B6H

S5018 W50140 0

1.2 0

0.8 0

0.04 0.03 4.50 . . . 0 0 6.00

B8

S5018 W50440 0

0.1 00.2 5 0.1 2

1.2 0

0.8 0

NIL(F)

K110 40

W21048

0.1 2

1.6 0

0.8 0

NI2(F)

K210 10

W22040

0.1 2

1.6 0

0.8 0

NI3

K313 10

W23040

0.1 2

1.6 0

0.8 0

NI4

K114 85

W21250

0.1 4

1.6 0

0.8 0

F1

K111

W21150

0.1

0.7

0.8

0.04 0.03 8.00 . . . 0 0 10.0 0 0.03 0.03 0.15 0.7 0 0 51.1 0 0.03 0.03 . . . 2.0 0 0 02.9 0 0.03 0.03 0.15 2.8 0 0 03.8 0 0.03 0.03 . . . 1.4 0 0 02.1 0 0.04 0.03 0.15 0.9

0.3 5

...

...

. . .

0.3 5

...

...

. . .

0.3 5

0.3 0

...

. . .

0.3 5

...

...

. . .

0.3 5

...

...

. . .

0.3 0

...

...

. . .

0.3 5

...

...

. . .

0.3 5

...

...

. . .

0.3 5

...

...

. . .

0.3 5

...

0.05(

. . .

...

0.3 5

...

...

. . .

...

0.3 5

...

...

. . .

0.3 5

0.3 5

...

...

. . .

0.5

0.3

...

...

.

0.4 00.6 5 0.4 00.6 5 0.4 00.6 5 0.9 01.2 0 0.4 00.6 5 0.9 01.2 0 0.4 00.6 5 0.4 00.6 5 0.8 01.2 0 0.3 5

G)

60

2

F2

K214 50

W20240

0.1 7

F3

K214 85

W21140

0.1 7

F4

K120 48

W20440

0.1 7

F5

K413 70

W22640

0.1 7

F6

K211 35

W21040

0.1 4

M1

...

W21240

0.1 0

M2

M3

M4

...

...

...

W21340

W22240

W22440

W

...

W21040

G

NOT SPECIFIED

0.1 0

0.1 0

0.1 0

0.1 2

01.5 0 1.2 52.2 5 1.2 52.2 5 1.6 0

0

0

0

0.8 0

0.04 0.03 . . . 0 0

0.8 0

0.04 0.03 . . . 0 0

0.8 0

0.04 0.03 0.60 0 0

1.2 01.8 0 0.8 01.8 5 0.6 01.6 0 0.9 01.8 0 0.9 01.8 0 1.3 02.2 5 0.5 01.6 0

0.8 0

0.03 0.03 0.65 0 0

0.8 0

0.03 0.03 0.65 0 0

0.8 0

0.04 0.03 0.15 0 0

0.8 0

0.04 0.03 0.35 0 0

0.8 0

0.03 0.03 0.65 0 0

0.8 0

0.03 0.03 0.80 0 0

0.8 0

0.04 0.03 0.45 0 0 0.70

01.7 0 0.4 00.8 0 0.7 01.1 0 0.4 00.8 0 2.0 02.8 0 1.5 02.2 5 1.2 52.0 0 1.4 02.1 0 1.8 02.6 0 2.0 02.8 0 0.4 00.8 0

5

5

. .

0.4 00.6 5 0.4 00.6 5 0.2 5

0.3 5

...

...

. . .

0.3 5

...

...

. . .

0.3 5

...

0.03(

. . .

0.3 00.8 0 0.6 0

0.5 0

...

...

. . .

0.4 0

...

...

. . .

0.3 5

0.3 0

...

0.03(

. . .

0.2 50.6 5 0.2 00.7 0 0.3 00.8 0 ...

0.3 0

...

0.3 0

...

0.3 0

...

0.3 00.7 5

...

G)

G)

0.03( G)

0.03( G)

. . . . . .

0.03( G)

...

. . .

Source: Ref 13

(A) THE WELD METAL SHALL BE ANALYZED FOR THE SPECIFIC ELEMENTS FOR WHICH VALUES ARE SHOWN IN THIS TABLE. IF THE PRESENCE OF OTHER ELEMENTS IS INDICATED, IN THE COURSE OF THIS WORK, THE AMOUNT OF THOSE ELEMENTS SHALL BE DETERMINED TO ENSURE THAT THEIR TOTAL (EXCLUDING IRON) DOES NOT EXCEED 0.50%. (B) SINGLE VALUES ARE MAXIMUM.

(C) THE LETTER "N" AS A SUFFIX TO CLASSIFICATION INDICATES THAT THE WELD METAL IS INTENDED FOR WELDS IN THE CORE BELT REGION OF NUCLEAR REACTOR VESSELS. THIS SUFFIX CHANGES THE LIMITS ON PHOSPHORUS, VANADIUM, AND COPPER AS FOLLOWS: P = 0.012% MAXIMUM; V = 0.05% MAXIMUM; CU = 0.08% MAXIMUM. "N" ELECRODES SHALL NOT BE COATED WITH COPPER OR ANY MATERIAL CONTAINING COPPER. THE F5 AND W CLASSIFICATIONS SHALL NOT BE DESIGNATED AS "N" WELD METALS. (D) A LOW-DILUTION AREA OF THE GROOVE WELD OR THE REDUCED SECTION OF THE FRACTURED TENSION TEST SPECIMEN MAY BE SUBSTITUTED FOR THE WELD PAD, AND SHALL MEET THE ABOVE REQUIREMENTS. IN CASE OF DISPUTE, THE WELD PAD SHALL BE THE REFEREE METHOD. (E) THE ELECRODE DESIGNATION FOR COMPOSITE ELECTRODES IS OBTAINED BY PLACING AN "EC" BEFORE THE APPROPRIATE WELD METAL CLASSIFICATION. (F) MANGANESE IN NI1 AND NI2 CLASSIFICATIONS MAY BE 1.80% MAXIMUM WHEN CARBON IS RESTRICTED TO 0.10% MAXIMUM. (G) A SINGLE VALUE SPANNING THE COLUMNS FOR VANADIUM, TITANIUM, AND ZIRCONIUM IS THE TOTAL OF THESE ELEMENTS.

FIG. 4 CLASSIFICATION SYSTEM FOR LOW-ALLOY STEEL ELECTRODES AND FLUXES USED IN SAW APPLICATIONS

F9PO-EB3-B3 is a complete designation for a flux-electrode combination. It refers to a flux that will produce weld metal

that, in the postweld heat-treated condition, will have a tensile strength of 620 to 760 MPa (90 to 110 ksi) and Charpy Vnotch impact strength of at least 27 J (20 ft · lbf) at - 18 °C (0 °F) when made with an EB3 electrode under the conditions called for in this specification. The composition of the weld metal will be B3 (Table 4). F9A2-ECM1-M1 is a complete designation for a flux when the trade name of the composite electrode used in the

classification tests is indicated as well. The designation refers to a flux that will produce weld metal that, with the M1

electrode, in the as-welded condition, will have a tensile strength of 620 to 760 MPa (90 to 110 ksi) and Charpy V-notch impact strength of at least 27 J (20 ft · lbf) at -29 °C (-20 °F) under the conditions called for in this specification. The composition of the weld metal will be M1 (Table 4). Stainless and Nickel-Base Alloys. For welding with corrosion-resistant chromium and chromium-nickel bare and

composite metal cored and stranded electrodes, the composition of the filler metal is classified according to AWS A5.9 (Table 5). There is no AWS classification of fluxes or of flux-electrode combinations for stainless steel SAW.

TABLE 5 COMPOSITION OF CORROSION-RESISTANT CHROMIUM AND CHROMIUM-NICKEL BARE AND COMPOSITE METAL CORED AND STRANDED ELECTRODES AWS ELECTRODE CLASSIFICATION

COMPOSITION, WT%(A)(B) C Cr Ni

ER209(D)

0.05

ER218

0.10

ER219

0.05

ER240

0.05

ER307

0.040.14 0.08

ER308(C) ER308H ER308L(C)

0.040.08 0.03

ER308MO

0.08

ER308MOL

0.04

ER309(D)

0.12

ER309L

0.03

ER310 ER312

0.080.15 0.15

ER316(C)

0.08

ER316H ER316L(D)

0.040.08 0.03

ER317

0.08

ER317L

0.03

ER318

0.08

20.524.0 16.018.0 19.021.5 17.019.0 19.522.0 19.522.0 19.522.0 19.522.0 18.021.0 18.021.0 23.025.0 23.025.0 25.028.0 28.032.0 18.020.0 18.020.0 18.020.0 18.520.5 18.520.5 18.0-

9.512.0 8.09.0 5.57.0 4.06.0 8.010.7 9.011.0 9.011.0 9.011.0 9.012.0 9.012.0 12.014.0 12.014.0 20.022.5 8.010.5 11.014.0 11.014.0 11.014.0 13.015.0 13.015.0 11.0-

Mo

(nb + ta)

1.53.0 0.75

...

0.75 0.75 0.51.5 0.75 0.75 0.75 2.03.0 2.03.0 0.75 0.75 0.75 0.75 2.03.0 2.03.0 2.03.0 3.04.0 3.04.0 2.0-

Mn

4.07.0 ... 7.09.0 ... 8.010.0 ... 10.513.5 ... 3.34.75 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 ... 1.02.5 8 × C 1.0-

Si

P

S

N

Cu

0.90

0.03

0.03

0.75

3.54.5 1.00

0.03

0.03

0.03

0.03

1.00

0.03

0.03

0.300.65 0.300.65 0.300.65 0.300.65 0.300.65 0.300.65 0.300.65 0.300.65 0.300.65 0.30065 0.300.65 0.300.65 0.300.65 0.300.65 0.300.65 0.30-

0.03

0.03

0.100.30 0.080.18 0.100.30 0.100.20 ...

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.75 0.75 0.75 0.75

20.0

14.0

3.0

2.5

0.65

2.5

0.60

0.03

...

3.0-4.0

1.52.0

0.15

0.015 0.020 . . .

3.0-4.0

1.02.5 ... 1.02.5 10 × C 1.02.5 MIN TO 1.0 MAX 1.0-1.4 1.02.5 ... 0.6

0.300.65 0.300.65 0.300.65

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.03

0.03

...

0.75

0.300.65 0.50

0.03

0.03

...

0.75

0.03

0.03

...

0.75

MIN TO 1.0 MAX 8 × C MIN TO 1.0 MAX 8 × C MIN TO 0.40 MAX ...

0.03

ER320

0.07

19.021.0

32.036.0

2.03.0

ER320LR(D)

0.025

19.021.0

32.036.0

2.03.0

ER321(E)

0.08 0.180.25 0.08

9.010.5 34.037.0 9.011.0

0.75

ER330

18.520.5 15.017.0 19.021.5

19.021.5 11.513.5 11.012.5 12.014.0 15.517.0 4.66.0 8.010.5 16.016.75 25.027.5 14.516.5

8.09.5 0.6

0.350.65 0.75

4.05.0 0.6

0.4.0.7 0.75

...

0.6

0.50

0.03

0.03

...

0.75

...

0.6

0.50

0.03

0.03

...

0.75

0.6

0.75

...

0.6

0.50

0.03

0.03

...

0.75

0.6

0.450.65 0.81.2 0.75

...

0.6

0.50

0.03

0.03

...

0.75

...

0.6

0.50

0.04

0.03

...

0.75

0.150.30 ...

0.250.75 0.40

0.75

0.04

0.03

...

0.40

0.02

0.02

0.015

3.254.00 0.20(G)

...

1.02.5

0.300.65

0.03

0.03

...

0.75

ER347(C)

ER349(F) ER410

0.070.13 0.12

ER410NIMO

0.06

ER420 ER430

0.250.40 0.10

ER502

0.10

ER505

0.10

ER630

0.05

ER26-1

0.01

ER16-8-2

0.10

0.5 4.55.0 (G)

7.59.5

0.75 0.75

0.751.50 1.02.0

Source: Ref 14

(A) ANALYSIS SHALL BE MADE FOR THE ELEMENTS FOR WHICH SPECIFIC VALUES ARE SHOWN IN THIS TABLE. IF, HOWEVER, THE PRESENCE OF OTHER ELEMENTS IS INDICATED IN THE COURSE OF ROUTINE ANALYSIS, FURTHER ANALYSIS SHALL BE MADE TO DETERMINE THAT THE TOTAL OF THESE OTHER ELEMENTS, EXCEPT IRON, IS NOT PRESENT IN EXCESS OF 0.50 %. (B) SINGLE VALUES SHOWN ARE MAXIMUM PERCENTAGES EXCEPT WHERE OTHERWISE SPECIFIED. (C) THESE GRADES ARE AVAILABLE IN HIGH-SILICON CLASSIFICATIONS THAT SHALL HAVE THE SAME CHEMICAL COMPOSITION REQUIREMENTS AS GIVEN BELOW WITH THE EXCEPTION THAT THE SILICON CONTENT SHALL BE 0.65 TO 1.00%. THESE HIGH-SILICON CLASSIFICATIONS SHALL BE DESIGNATED BY THE ADDITION `SI' TO THE STANDARD

CLASSIFICATION DESIGNATIONS INDICATED BELOW. THE FABRICATOR SHOULD CONSIDER CAREFULLY THE USE OF HIGH-SILICON FILLER METALS IN HIGHLY RESTRAINED FULLY AUSTENITIC WELDS. A DISCUSSION OF THE PROBLEM IS PRESENTED IN REF 14. (D) CARBON SHALL BE REPORTED TO THE NEAREST 0.01% EXCEPT FOR THE CLASSIFICATION E320LR, FOR WHICH CARBON SHALL BE REPORTED TO THE NEAREST 0.005%. (E) TITANIUM: 9 × C MIN TO 1.0 MAX. (F) TITANIUM: 0.10 TO 0.30. TUNGSTEN IS 1.25 TO 1.75%. (G) NICKEL, MAX: 0.5 MINUS THE COPPER CONTENT. Like stainless steels, nickel and nickel alloys are classified according to the composition of the filler metal only (Ref 15). Classifications of filler metal are shown in Table 6. There is no AWS classification of fluxes or of flux-electrode combinations for nickel alloy SAW.

TABLE 6 COMPOSITION OF NICKEL AND NICKEL-ALLOY BARE ELECTRODES AND RODS

AWS CLASSIFICATIO N

UNS NO.

COMPOSITION, WT%(A)(B) C

MN

FE

P

S

SI

CU

N(C)

CO

AL

TI

CR

(NB + TA)

MO

V

W

OTHER ELEMENT S, TOTAL

ERNI-1

N0206 1

0.15

1.0

1.0

0.03

0.01 5

0.7 5

0.25

...

1.5

2.03.5

...

...

...

...

...

0.50

ERNICU-7

N0406 0

0.15

4.0

2.5

0.02

0.01 5

1.2 5

BA L

...

1.25

1.53.0

...

...

...

...

...

0.50

ERNICR-3

N0608 2

0.10

2.5- 3.0 3.5

0.03

0.01 5

0.5 0

0.50

93.0 MI N 62.0 69.0 67.0 MI N 70.0 MI N 67.0 MI N 38.0 46.0 50.0 55.0 BA L

(D)

...

0.75

2.0- . . . 3.0(E

...

...

0.50

...

...

0.50

ERNLCRFE-5

N0606 2

0.08

1.0

6.010.0

0.03

0.01 5

0.3 5

...

...

2.03.5

...

0.20

0.60 -1.2

18.0 22.0 14.0 17.0 14.0 17.0 19.5 23.5

...

2.5

0.20 0.80 ...

0.65 1.15 ...

17.0 21.0 1.0

BA L

0.20

...

...

6.08.0

0.50

BA L

2.5

...

...

4.06.0

...

0.50

BA

1.0

...

...

1.0

...

0.50

ERNLCRFE-6

N0709 2

0.08

2.0- 8.0 2.7

0.03

0.01 5

0.3 5

0.50

ERNIFECR-1

N0806 5

0.05

1.0

0.03

0.03

0.5 0

1.50 -3.0

ERNIFECR-2(F)

N0771 8

0.08

0.3 5

22.0 MI N BA L

0.01 5

0.01 5

0.3 5

0.30

ERNIMO-1

N1000 1

0.08

1.0

4.07.0

0.02 5

0.03

1.0

0.50

ERNIMO-2

N1000 3

0.04- 1.0 0.08

5.0

0.01 5

0.02

1.0

0.50

ERNIMO-3

N1000 4

0.12

1.0

4.07.0

0.04

0.03

1.0

ERNIMO-7

N1066

0.02

1.0

2.0

0.04

0.03

0.1

(D)

...

...

)

1.5- . . . 3.0(E )

...

...

...

...

0.50

...

2.53.5

...

...

0.50

4.75 5.50 ...

2.80 3.30 26.0 30.0 15.0 18.0 23.0 26.0 26.0

...

...

0.50

0.20 0.40 0.50

1.0

0.50

0.50

0.50

0.60

1.0

0.50

...

1.0

0.50

5 ERNICRMO-1

N0600 7

ERNICRMO-2

N0600 2

ERNICRMO-3

N0662 5

ERNICRMO-4

0 0.05

L

1.0- 18.0 2.0 21.0 0.05- 1.0 17.0 0.15 20.0 0.10 0.5 5.0 0

0.04

0.03

1.0

1.52.5

BA L

2.5

...

...

0.04

0.03

1.0

0.50

BA L

0.50 -2.5

...

...

0.02

0.01 5

0.5 0

0.50

...

0.40

0.40

N1027 6

0.02

1.0

4.07.0

0.04

0.03

0.0 8

0.50

58.0 MI N BA L

2.5

...

...

ERNICRMO-7

N0645 5

0.01 5

1.0

3.0

0.04

0.03

0.0 8

0.50

BA L

2.0

...

0.70

ERNICRMO-8

N0697 5

0.03

1.0

BA L

0.03

0.03

1.0

0.71.20

...

...

ERNICRMO-9

N0698 5

0.01 5

1.0

0.04

0.03

1.0

1.52.5

5.0

...

0.70 1.50 ...

ERNICRMO-10

N0602 2

0.01 5

0.5 0

18.0 21.0 2.06.0

47.0 52.0 BA L

0.02

0.01 0

0.0 8

0.50

BA L

2.5

...

...

ERNICRMO-11

N0603 0

0.03

1.5

0.04

0.02

0.8 0

1.02.4

BA L

5.0

...

...

ERNICRCOMO-1

N0661 7

0.05- 1.0 0.15

13.0 17.0 3.0

0.03

0.01 5

1.0

0.50

BA L

10.0 15.0

0.80 1.50

0.60

21.0 23.5 20.5 23.0 20.0 23.0 14.5 16.5 14.0 18.0 23.0 26.0 21.0 23.5 20.0 22.5 28.0 31.5 20.0 24.0

1.75 2.50 ...

30.0 5.57.5

...

1.0

0.50

8.010.0

...

0.20 -1.0

0.50

8.010.0

...

...

0.50

15.0 17.0 14.0 18.0 5.07.0

0.35

3.04.5

0.50

...

0.50

0.50

...

...

0.50

0.50

6.08.0

...

1.5

0.50

...

12.5 14.5 4.06.0

0.35

2.54.5

0.50

...

1.54.0

0.50

8.010.0

...

...

0.50

3.15 4.15 ...

...

...

0.30 1.50 ...

Source: Ref 15

(A) THE FILLER METAL SHALL BE ANALYZED FOR THE SPECIFIC ELEMENTS FOR WHICH VALUES AVE SHOWN IN THIS TABLE. IN THE COURSE OF THIS WORK, IF THE PRESENCE OF OTHER ELEMENTS IS INDICATED, THE AMOUNT OF THOSE ELEMENTS SHALL HE DETERMINED TO ENSURE THAT THEIR TOTAL DOES NOT EXCEED THE LIMIT SPECIFIED FOR "OTHER ELEMENTS, TOTAL" IN

THE LAST COLUMN OF THE TABLE. (B) SINGLE VALUES ARE MAXIMUM, EXCEPT WHERE OTHERWISE SPECIFIED. (C) INCLUDES INCIDENTAL COBALT. (D) COBALT: 0.12 MAXIMUM, WHEN SPECIFIED. (E) TANTALUM: 0.30 MAXIMUM, WHEN SPECIFIED. (F) BORON IS 0.006% MAXIMUM.

Submerged Arc Welding Jonathan S. Ogborn, The Lincoln Electric Company

Personnel Considerations Employee Training. Submerged arc welding is an automatic or semiautomatic process carried out primarily in the flat or

horizontal position with the arc obscured from view and does not require a high level of skill. The primary expenditures both in cost and time are consumed in designing jigs and fixtures to hold and position the workpiece so that SAW can be performed. While the actual welding process does not require a high degree of skill, the setup procedure requires skilled personnel with an aptitude for fabricating the fixtures and clamps necessary to position the workpiece in an acceptable position (usually flat). Health and Safety in the Workplace. The main safety concerns in welding are electric shock, fume inhalation, and

burns from both the arc ultraviolet source and the hot-metal infrared source (Ref 16). To reduce the risk of electrical shock, all equipment must be properly grounded and welding cables must be in good condition. The fume is nearly eliminated in SAW. However, certain vaporized elements (for example, chromium, cobalt, manganese, nickel, and vanadium) can be potentially dangerous. To prevent exposure to these elements, the welding area should have adequate ventilation. In SAW, the arcflash and its accompanying spatter are nearly eliminated because of the slag and the depth of the flux pile. However, eye protection should always be worn. Additional information is available in the article "Safe Practices" in this Volume.


16. SAFETY IN WELDING AND CUTTING, AWS, 1988 Submerged Arc Welding Jonathan S. Ogborn, The Lincoln Electric Company

SAW Parameters While SAW is the most inexpensive and efficient process for making large, long, and repetitive welds, much time and energy are required to prepare the joint. Care must be taken to line up all joints to have a consistent gap in groove welds and to provide backing plates and flux dams to prevent spillage of flux and molten metal. Once all the pieces are clamped or tacked in place, welding procedures and specifications should be consulted before welding begins. Procedural Variations and Effect on Weld Bead Characteristics Procedural variations in SAW include current, voltage, electrical stickout (distance from last electrical contact to plate), travel speed, and flux depth. Variation in any of these parameters will affect the shape and penetration of the weld, as well as the integrity of the weld deposit. Weld Current. Because the welding current controls such parameters as deposition rate, penetration, and dilution, it is

the most important welding variable. An increase in welding current at a constant voltage will decrease the flux-to-wire ratio, while a decrease in current will increase the flux-to-wire ratio. The effect of current variation on weld bead profile is shown in Fig. 5. Welds made at excessively low current will tend to have little penetration and higher width-to-depth ratios. Welds made at an excessively high current will have deep penetration, high dilution, more shrinkage, and excess buildup. Low current will also produce a less stable arc than higher currents.

FIG. 5 EFFECT OF VARIATION IN WELDING CURRENT ON WELD BEAD PROFILE. (A) EXCESSIVELY LOW CURRENT. (B) EXCESSIVELY HIGH CURRENT. (C) RECOMMENDED CURRENT

The direction of current flow will also affect the weld bead profile. The current may be direct with the electrode positive (reverse polarity), electrode negative (straight polarity), or alternating. Reverse polarity is most commonly used. For a given set of welding conditions, reverse polarity will produce wider beads with more penetration at a lower deposition rate than straight polarity. Straight-polarity welding will contribute to narrower beads with less penetration and more buildup. Because straight polarity reduces base plate dilution, it is frequently used in surfacing applications. For the same welding current, the deposition rate with straight polarity is higher than with reverse polarity. Straight polarity is preferred for poor fitup. The bead shape, penetration, and deposition rate for alternating current fall between those of straight and reverse polarity. Alternating Current is used when welding current exceeds 1000 A and on multiple-wire applications to reduce arc blow and arc interaction. In SAW, the current density in the electrode also plays a role in bead shape and penetration. Smaller-diameter electrodes with a high current density will produce narrower beads with deeper penetration than larger-diameter electrodes. Largerdiameter electrodes are able to bridge larger root openings. In cases where a given current can be achieved with two different electrode diameters, the smaller electrode will produce the higher deposition rate. Weld Voltage. Like current, welding voltage will affect the bead shape and the weld deposit composition. Increasing the

arc voltage at a constant current will increase the flux-to-wire electrode ratio, while decreasing the voltage will reduce the flux-to-electrode ratio. The effect of the magnitude of arc voltage on bead shape is shown in Fig. 6. Increasing the arc voltage will produce a longer arc length and a correspondingly wider, flatter bead with less penetration. Higher voltage will increase flux consumption, which could then change deposit composition and properties. Slightly increasing the arc voltage will help the weld to bridge gaps when welding in grooves. Excessively high voltage will produce a hat-shaped concave weld, which has low resistance to cracking and a tendency to undercut. Lower voltages will shorten the arc length and increase penetration. Excessively low voltage will produce an unstable arc and a crowned bead, which has an uneven contour where it meets the plate.

FIG. 6 EFFECT OF VARIATION IN WELDING VOLTAGE AT CONSTANT CURRENT ON WELD BEAD PROFILE. (A) EXCESSIVELY LOW VOLTAGE. (B) EXCESSIVELY HIGH VOLTAGE

Electrical Stickout. In SAW, the current flowing in the electrode between the contact tube and the arc (electrode

extension) will cause some I2R heating, resulting in a voltage drop across that length of electrode. This resistance heating and subsequent voltage drop can be used to obtain higher deposition rates. Normal electrode extension for solid SAW wire is approximately 8 to 12 times the electrode diameter. As this length increases at a constant current, so does the resistance heating and the melt-off rate. To compensate for the voltage drop and the increase in wire-feed speed, the voltage must be increased to obtain a properly shaped bead. Extending the electrode 20 to 40 times the diameter can

increase deposition rates by more than 50% (Ref 17). Although higher deposition rates can be achieved by extending the electrode, keeping the electrode aligned with the joint becomes increasingly difficult as the extension increases. Travel Speed. Variations in travel speed at a set current and voltage also affect bead shape. As welding speed is

decreased, heat input per length of joint increases, and the penetration and bead width increase. The penetration will increase until molten metal begins to flow under the arc and interfere with heat flow at excessively slow speeds. Excessively high travel speeds will promote a crowned bead as well as the tendency for undercut and porosity. Flux layer depth is another variable that will alter the appearance, penetration, and quality of a submerged arc weld. If

the flux layer is too deep, a greater-than-normal amount of flux will be melted, resulting in weld beads that are narrower than normal. Some surface imperfections may also appear because gases may be trapped by the deep flux layer. If the flux layer is too shallow, the arc will flash through, and the bead will have a rough appearance or porosity due to lack of shielding from the atmosphere. The correct depth of flux is just enough to prevent flash through. This will allow welding gases to escape while providing adequate protection. Sources of Defects in SAW The fact that SAW is a high heat input process under a protective blanket of flux greatly decreases the chance of weld defects. However, defects such as lack of fusion, slag entrapment, solidification cracking, hydrogen cracking, or porosity occasionally occur. Insufficient Fusion and Slag Entrapment. Lack-of-fusion defects and slag entrapment are most commonly caused by improper bead placement or procedure. Improper placement can cause the weld metal to roll over and trap slag underneath, or if the weld bead is placed away from the edge to be joined, the liquid metal may not fuse to the base material. A crown-shaped bead caused by low welding voltage may also contribute to slag entrapment and lack of fusion by not allowing the liquid metal to spread out evenly. Solidification cracking of SAW along the center of the bead is usually due to bead shape, joint design, or incorrect choice of welding consumables. A convex bead shape with a bead width-to-depth ratio greater than one will decrease solidification cracking tendencies. If weld penetration is too deep, the shrinkage stresses may cause centerline cracking. Joint design may also contribute to excessive shrinkage stresses, again increasing the risk of solidification cracking. Because cracking is related to stresses in the weld, high-strength materials will have a greater tendency to crack. Therefore, special care must be taken to generate proper bead shape, preheat temperatures, and interpass temperatures, in addition to correct electrode and flux combinations, when welding these materials. Hydrogen Cracking. Unlike solidification cracking, which appears immediately after welding, hydrogen cracking is a

delayed process and may occur from several hours to several days after welding has been completed. To minimize hydrogen cracking, all possible sources of hydrogen (for example, water, oil, grease, and dirt) present in the flux, electrode, or joint should be eliminated. The flux, electrode, and plate should be clean and dry. To prevent moisture pickup, fluxes and electrodes should be stored in moisture-resistant containers in dry areas. If a flux or electrode becomes contaminated with moisture, it should be dried according to manufacturer recommendations. Selection of the consumable, especially when welding high-strength steels that are more susceptible to hydrogen cracking, can also play a role in hydrogen cracking. Special consumables have recently been developed that produce weld deposits that are very low in diffusible hydrogen. To further reduce hydrogen-related cracking, the joint to be welded should be preheated. Because hydrogen is fairly mobile in steel at temperatures above 95 °C (200 °F), the recommended preheat temperatures should be followed to allow most of the hydrogen to escape and to reduce the risk of hydrogen damage (Ref 18). In thick weldments, maintaining preheat for several hours after welding is completed will also reduce the risk of hydrogen cracking. Porosity caused by trapped gas is uncommon in SAW because of the protection provided by the flux. When porosity

does occur, it may be in the form of internal porosity or as depressions on the weld bead surface. The gas bubbles that cause porosity originate either from a lack of protection from the atmosphere or from contaminants such as water, oil, grease, and dirt. To reduce porosity in SAW, the weld should have sufficient flux coverage, and all water, grease, and dirt should be removed from the plate, electrode, and flux. Another cause of porosity in SAW is excessive travel speed. Travel

at excessively high speeds will not allow the gas bubbles to escape from the weld, and the bubbles may become trapped in the weld metal at the slag-to-metal interface.


17. HOW TO MAKE SINGLE ELECTRODE SUBMERGED ARC WELDS, PUBLICATION S604, THE LINCOLN ELECTRIC CO., 1991, P 11 18. F.R. COE, WELD. WORLD, VOL 14 (NO. 1/2), 1976, P 1-7 Submerged Arc Welding Jonathan S. Ogborn, The Lincoln Electric Company

References

1. R.L. O'BRIEN, WELDING HANDBOOK, VOL 11, AWS, 1991, P 191-232 2. D.L. OLSON ET AL., SUBMERGED ARC WELDING, VOL 6, 9TH ED., METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 114-152 3. THE PROCEDURE HANDBOOK OF ARC WELDING, THE LINCOLN ELECTRIC CO., 1973, P 6.3-1 TO 6.3-24 4. "SPECIFICATION FOR CARBON STEEL ELECTRODES AND FLUXES FOR SUBMERGED ARC WELDING," A5.17-89, AWS, 1989, P 18 5. THE PROCEDURE HANDBOOK OF ARC WELDING, THE LINCOLN ELECTRIC CO., 1973, P 3.2-3, 6.3-21 TO 6.3-22, 6.3-49 TO 6.3-58 6. THE PROCEDURE HANDBOOK OF ARC WELDING, THE LINCOLN ELECTRIC CO., 1973, P 6.3-26 TO 6.3-73 7. R.S. BROWN ET AL., ARC WELDING OF STAINLESS STEELS, VOL 6, 9TH ED., METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 327-329 8. K.C. ANTONY ET AL., HARDFACING, VOL 6, 9TH ED., METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 784 9. D.R. THOMAS ET AL., WELD OVERLAYS, VOL 6, 9TH ED., METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 804-819 10. "SPECIFICATION FOR CARBON STEEL ELECTRODES AND FLUXES FOR SUBMERGED ARC WELDING," A5.17-89, AWS, 1989, P 17-18 11. S.S. TULIANI, T. BONISZEWSKI, AND N.F. EATON, NOTCH TOUGHNESS OF COMMERCIAL SUBMERGED ARC WELD METAL, WELD. MET. FABR., VOL 8, 1969, P 327-339 12. "SPECIFICATION FOR CARBON STEEL ELECTRODES AND FLUXES FOR SUBMERGED ARC WELDING," A5.17-89, AWS, 1989, P 2-3 13. "SPECIFICATION FOR LOW ALLOY STEEL ELECTRODES AND FLUXES FOR SUBMERGED ARC WELDING," A5.23-90, AWS, 1990, P 2-6 14. "SPECIFICATION FOR CORROSION RESISTING CHROMIUM AND CHROMIUM-NICKEL STEEL BARE AND COMPOSITE METAL CORED AND STRANDED WELDING ELECTRODES AND WELDING RODS," A5.9-81, AWS, 1981, P 3, 4 15. "SPECIFICATION FOR NICKEL AND NICKEL ALLOY BARE WELDING ELECTRODES AND RODS," A5.14-89, AWS, 1989, P 2, 3 16. SAFETY IN WELDING AND CUTTING, AWS, 1988 17. HOW TO MAKE SINGLE ELECTRODE SUBMERGED ARC WELDS, PUBLICATION S604, THE LINCOLN ELECTRIC CO., 1991, P 11 18. F.R. COE, WELD. WORLD, VOL 14 (NO. 1/2), 1976, P 1-7

Stud Arc Welding Harry A. Chambers, TRW Nelson Stud Welding Division

Introduction STUD ARC WELDING (SW), also known as arc stud welding, is a commonly used method for joining a metal stud, or fastener, to a metal workpiece. The process has been used as an alternative metal-fastening method since the 1940s. Millions of specially designed and manufactured metal studs are welded by this process every week in such diverse industries as construction, shipbuilding, automotive, and hard goods, as well as in miscellaneous industrial applications. This article will serve as a basic information source for those interested in accomplishing one-sided, no-hole attachment of metal fasteners. The SW process represents an alternative to other welding processes, and is also a substitute for other fastening procedures, such as drilling and tapping, bolting, and self-tapping screws. Stud Arc Welding Harry A. Chambers, TRW Nelson Stud Welding Division

Process Overview Stud arc welding is similar to many other welding processes, including arc and percussion welding, in that the base (weld end) of a specifically designed stud is joined to a base material by heating both parts with an arc that is drawn between the two. Equipment that is unique to this process regulates the arc length and arc dwell time. After an arc is struck, the stud weld end and the workpiece surface are brought to the proper temperature for joining and, after a controlled period of time, the two heated surfaces are brought together under pressure, creating a metallurgical bond capable of developing the full strength of the stud. There are two basic types of stud arc welding, which are differentiated by the source of welding power. One type uses direct current (dc) power provided by a transformer/rectifier or a motor generator similar to that used in the shielded metal arc welding (SMAW) process. The second type uses power discharged from a capacitor storage bank. The process based on a dc power source is known as stud arc welding, whereas the process that utilizes capacitors is known as capacitor discharge stud welding (CDSW). Both the SW and CDSW processes overlap in some areas of application. Generally, the SW process is used in applications that require similar stud and workpiece metals, the workpiece thickness is greater in relation to the stud diameter, and an accommodation must be made for the stud flash (fillet). The term "flash" is preferred to the term "fillet," because the metal that forms the flash during the stud arc welding process is expelled, rather than added, as occurs with other welding processes. In contrast, the CDSW process is used extensively when welding to thin sheet metal, and is used frequently with dissimilar workpiece and stud alloys. It is also used in cases where marks on the opposite side of the workpiece must be avoided or minimized. With this process, the stud diameter is limited to smaller sizes. The CDSW process is more fully described in the article "Capacitor Discharge Stud Welding" in this Volume. The factors on which process selection should be based are fastener size, base-metal thickness, base-metal composition, and reverse-side marking requirements (Table 1).

TABLE 1 STUD-WELDING PROCESS SELECTION PARAMETERS

STUD SHAPE ROUND SQUARE

STUD ARC WELDING

CAPACITOR DISCHARGE STUD WELDING GAP AND DRAWN ARC METHOD CONTACT METHODS

A A

A A

A A

RECTANGULAR IRREGULAR STUD DIAMETER OR AREA

A A

A A

A A

1 1 TO IN.) 16 8 1 1 3.2 TO 6.4 MM ( TO IN.) 8 4 1 1 6.4 TO 12.7 MM ( TO IN.) 4 2 1 12.7 TO 25.4 MM ( TO 1 IN.) 2

D

A

A

C

A

A

A

B

B

A

D

D

UP TO 32.3 MM2 (0.05 IN.2) OVER 32.3 MM2 (0.05 IN.2) STUD METAL CARBON STEEL STAINLESS STEEL ALLOY STEEL ALUMINUM BRASS BASE METAL CARBON STEEL STAINLESS STEEL ALLOY STEEL ALUMINUM BRASS BASE-METAL THICKNESS UNDER 0.4 MM (0.015 IN.) 0.4 TO 1.6 MM (0.015 TO 0.062 IN.) 1.6 TO 3.2 MM (0.062 TO 0.125 IN.) OVER 3.2 MM (0.125 IN.) STRENGTH CRITERIA HEAT EFFECT ON EXPOSED SURFACES WELD FILLET CLEARANCE STRENGTH OF STUD GOVERNS STRENGTH OF BASE METAL GOVERNS

C A

A D

A D

A A B B C

A A C A A

A A C A D

A A B B C

A A A A A

A A C A D

D C B A

A A A A

B A A A

B B A A

A A A A

A A A A

1.6 TO 3.2 MM (

A. applicable without special procedures or equipment; B. applicable with special techniques or for specific applications that justify preliminary trials or testing to develop welding procedure and technique; C, limited application; D, not recommended-welding methods not developed at this time. Source: Ref 1 Equipment. As shown in Fig. 1, the basic equipment used for stud arc welding consists of a control system, which

regulates the arc time and controls gun movement; a fixed or portable stud-welding gun, which holds the stud in position during the welding process to create the proper arc length and joining pressure; and connecting cables, which must be connected to a separate source of dc power. The other items that are needed to weld the workpiece are the studs themselves and ceramic arc shields, or ferrules. The equipment used for stud welding is comparable, in terms of size, portability, and ease of operation, to the equipment used in the SMAW process. Although the initial cost of stud-welding systems varies with the method selected, stud size, and productivity requirements, it is generally competitive with other fastening methods on an in-place or finished part cost basis.

FIG. 1 STUD ARC WELDING CONTROL SYSTEM. THE CONTROL SYSTEM MUST BE CONNECTED TO A DC POWER SOURCE FOR WELDING. IN MOST APPLICATIONS, THE STUD (ELECTRODE) SHOULD BE NEGATIVE.

A stud-welding control system that has been integrated in a transformer/rectifier power source is shown in Fig. 2. This type of equipment, which is the most widely used, is called a power/control system. It can weld studs with diameters up to 28.5 mm (1

1 in.). Either a single gun or dual guns can be used. Although a light-duty control system can weigh 11.3 kg 8

(25 lb), the system shown in Fig. 2 weighs approximately 450 kg (1000 lb) and can be put on a wheeled cart for mobility.

FIG. 2 TYPICAL INTEGRATED POWER/CONTROL SYSTEM FOR STUD ARC WELDING

A typical stud-welding gun (Fig. 3) is usually available in at least two sizes: standard duty, for use with studs up to 16 mm (

5 in.) in diameter, and heavy duty, for use on larger-diameter studs. The weight of the gun can vary from approximately 8

1.8 kg (4 lb), for a standard-duty gun, to 6.8 kg (15 lb), for heavy-duty models.

FIG. 3 TYPICAL STANDARD-DUTY GUN FOR STUD ARC WELDING

Operation. The stud arc welding process utilizes the same principles as any other arc welding procedure. First, the stud,

which acts as an electrode, is inserted into a chuck on the end of the gun, surrounded by a ceramic ferrule, and positioned against the workpiece (Fig. 4a). Next, the gun trigger is depressed, which starts the automatic weld cycle by energizing a solenoid coil within the gun body that lifts the stud off the work and draws an arc. The arc melts the end of the stud and a portion of the workpiece (Fig. 4b). After a preset arc time (set on the control unit), the welding current is shut off and the solenoid is de-energized (Fig. 4c). A mainspring in the gun forces the stud into the molten pool of metal, producing a fullstrength weld, which is shown in Fig. 4(d) after the gun has been lifted off the stud and the ceramic ferrule removed. The result is a full-penetration, full-strength stud-to-workpiece weld, as shown in Fig. 5. This weld is similar to that obtained with other types of arc-welding processes.

FIG. 4 STUD ARC WELDING PROCESS. (A) GUN IS PROPERLY POSITIONED. (B) TRIGGER IS DEPRESSED AND STUD IS LIFTED, CREATING AN ARC. (C) ARCING PERIOD IS COMPLETED AND STUD IS PLUNGED INTO MOLTEN POOL OF METAL ON BASE MATERIAL. (D) GUN IS WITHDRAWN FROM WELDED STUD AND FERRULE IS REMOVED.

FIG. 5 MACROSECTION OF LOW-CARBON STEEL STUD WELD

Studs and Ferrules. With the exception of special processes, the stud arc welding of steel, stainless steel, and aluminum

studs requires the use of a specifically designed weld stud and ceramic ferrule. Studs. A wide range of studs are made by stud-welding manufacturers for the SW process. They vary in size and weld-

base configuration, depending on the application requirements. The typical stud styles that are available include threaded, unthreaded, headed, and rectangular studs (Fig. 6). The steel stud diameters that are commonly welded using the SW process range from 2.7 to 25.4 mm (0.105 to 1.00 in.).

FIG. 6 COMMON STUD CONFIGURATIONS FOR STUD ARC WELDING

Full-strength welds result when using studs made from low-carbon steel that conforms to ASTM A-108 grades C-1010 through C-1020, although other steel grades can be used on a special-application qualified basis (Ref 2). Stainless steel studs are typically made from the austenitic series, including AISI 302, 302 HQ, 304, 305, 308, 309, 310, 316, 321, and 347, in both normal-carbon and low-carbon varieties. Stainless steel 303 is not an acceptable stud alloy. Aluminum studs are usually made from aluminum-magnesium alloys, including 5183, 5356, 5556, 5086, and 5456 (Ref 3). Steel and stainless steel arc welded studs that are larger than 6.4 mm (0.250 in.) in diameter use a welding flux to stabilize the welding arc and to deoxidize the weld area. The most commonly used flux is commercially pure aluminum, which is installed onto the weld end of the stud by thermal spraying, staking, or pressing a slug into a drilled hole in the end of the stud (Fig. 7). Studs that are 6.4 mm (0.250 in.) and under usually do not require fluxing.

FIG. 7 METHODS OF FLUX LOADING STUDS. (A) PRESSED IN SLUG. (B) STAKED ON WASHER. (C) COATING

Aluminum studs do not have to be flux loaded. Instead, the weld end of the stud has a conical or cylindrical tip that is shaped into the parent metal to help initiate the arc and control the arc length. To prevent the oxidation and embrittlement of the stud/plate weld zone during aluminum stud arc welding, the weld area is purged with an inert gas during the weld cycle. Argon gas is commonly used for this purpose, although helium also can be used. The typical gas flow rates for aluminum studs of various diameters are shown in Table 2.

TABLE 2 GAS FLOW RATES FOR ALUMINUM STUD ARC WELDING

STUD WELD BASE DIAMETER SHIELDING GAS FLOW(A) mm in. l/min ft3/h 1 6.4 7.1 15 7.9 9.5 11.1 12.7 (A)

4 5 16 3 8 7 16 1 2

7.1

15

9.4

20

9.4

20

9.4

20

SHIELDING GAS, 99.95% PURE ARGON

When stud arc welding low-carbon and stainless steel studs, the stud (electrode) should be negative. Aluminum stud arc welding should be set up so that the stud is positive. Obtaining a full-quality stud arc weld requires a sufficient amount of total energy input to the weld joint in order to produce melting and complete metallurgical bonding of the stud and workpiece. The energy input or weld current that is necessary depends on the stud diameter. Other parameters in the stud arc welding process are: arc voltage, arc time, and plunge. Arc voltage is a function of the arc length, which is set as "lift," or the length the stud is drawn away from the workpiece during weld-cycle initiation. Plunge is the length of stud that extends past the end of the ceramic ferrule and is available for melt off during the weld cycle. Studs are reduced in length during the weld cycle. Table 3 shows the typical reductions for steel and aluminum studs as a result of welding.

TABLE 3 TYPICAL STUD WELDING SETTINGS STUD BASE DIAMETE R mm in.

LENGTH REDUCTI ON mm

in.

DOWNHAND WELDING CURRE TIME LIFT NT, A s cycl m in. es m

STEEL AND STAINLESS STEEL 4.8 0.1 3. 0.12 300 0.1 87 2 5 5 6.4 0.2 3. 0.12 450 0.1 50 2 5 7 7.9 0.3 3. 0.12 500 0.2 12 2 5 5 9.5 0.3 3. 0.12 550 0.3 75 2 5 3 11. 0.4 3. 0.12 675 0.4 1 37 2 5 2 12. 0.5 4. 0.12 800 0.5 7 00 8 5 5 15. 0.6 4. 0.18 1200 0.6 9 25 8 7 7 19. 0.7 4. 0.18 1500 0.8 1 50 8 7 4 22. 0.8 4. 0.18 1700 1.0 2 75 8 7 0 25. 1.0 6. 0.25 1900 1.4 4 00 4 0 0 (A) ALUMINUM 4.8 0.1 3. 0.12 150 0.2 87 2 5 5 6.4 0.2 3. 0.12 200 0.4 50 2 5 0 7.9 0.3 3. 0.12 250 0.5 12 2 5 0 9.5 0.3 3. 0.12 325 0.6 75 2 5 5

8-9 1012 15 20 25 33 40 5055 6065 85

15 24 30 39

PLUNGE

m m

in.

OVERHEAD WELDING CURRE TIME LIFT NT, A s cycl m in. es m

1. 6 1. 6 1. 6 1. 6 1. 6 1. 6 2. 4 2. 4 3. 2 3. 2

0.0 62 0.0 62 0.0 62 0.0 62 0.0 62 0.0 62 0.0 93 0.0 93 0.1 25 0.1 25

3. 2 3. 2 3. 2 3. 2 3. 2 3. 2 4. 7 4. 7 6. 4 6. 4

0.1 25 0.1 25 0.1 25 0.1 25 0.1 25 0.1 25 0.1 87 0.1 87 0.2 50 0.2 50

300

2. 4 2. 4 2. 4 3. 2

0.0 93 0.0 93 0.0 93 0.1 25

3. 2 3. 2 3. 2 3. 2

0.1 25 0.1 25 0.1 25 0.1 25

150

450 500 550 675 800 1200 1500 1700 2050

200 250 325

0.1 5 0.1 5 0.2 5 0.3 3 0.4 2 0.5 5 0.6 7 0.8 4 1.0 0 1.4 0

8-9

0.2 5 0.4 0 0.5 0 0.6 5

15

1012 15 20 25 33 40 5055 65 72

24 30 39

PLUNGE

m m

in.

VERTICAL WELDING CURRE TIME LIFT NT, A s cycl m in. es m

1. 6 1. 6 1. 6 1. 6 1. 6 1. 6 2. 4 2. 4 3. 2 3. 2

0.0 62 0.0 62 0.0 62 0.0 62 0.0 62 0.0 62 0.0 93 0.0 93 0.1 25 0.1 25

3. 2 3. 2 3. 2 3. 2 3. 2 3. 2 4. 7 4. 7 6. 4 6. 4

0.1 25 0.1 25 0.1 25 0.1 25 0.1 25 0.1 25 0.1 87 0.1 87 0.2 50 0.2 50

300

0.1 8-9 1. 0.0 5 6 62 450 0.1 10- 1. 0.0 7 12 6 62 500 0.2 15 1. 0.0 5 6 62 600 0.3 20 1. 0.0 3 6 62 750 0.3 20 1. 0.0 3 6 62 875 0.4 28 1. 0.0 7 6 62 1275 0.6 36 1. 0.0 0 6 62 1700 0.7 50 2. 0.0 3 4 93 NOT RECOMMENDED

2. 4 2. 4 2. 4 2. 4

0.0 93 0.0 93 0.0 93 0.0 93

3. 2 3. 2 3. 2 3. 2

0.1 25 0.1 25 0.1 25 0.1 25

180

PLUNGE

m m

in.

3. 2 3. 2 3. 2 3. 2 3. 2 3. 2 4. 7 4. 7

0.125

3. 2 3. 2 3. 2 3. 2

0.125

0.125 0.125 0.125 0.125 0.125 0.187 (A)

0.187 (A)

NOT RECOMMENDED

225 275 350

0.2 0 0.3 0 0.4 0 0.6 0

12 18 24 36

2. 4 2. 4 2. 4 3. 2

0.0 93 0.0 93 0.0 93 0.1 25

0.125 0.125 0.125

11. 1 12. 7

0.4 37 0.5 00

3. 2 3. 2

0.12 400 5 0.12 460 5

0.8 0 0.9 0

48 54

3. 2 3. 2

0.1 25 0.1 25

3. 9 4. 7

0.1 56 0.1 87

400 460

0.8 0 0.9 0

48 54

3. 9 4. 7

0.1 56 0.1 87

3. 9 4. 7

0.1 56 0.1 87

(A) STUD ARC WELDED ALUMINUM STUDS REQUIRE INERT GAS SHIELDING: SEE TABLE 2.

430 475

0.7 0 0.8 0

42 48

3. 2 3. 2

0.1 25 0.1 25

3. 9 4. 7

0.156 0.187

There is a range of setting combinations for all stud diameters. The same total energy input can be obtained by varying current input and arc time. For example, a low current input can be compensated for, to some extent, by increasing weld time. Table 3 also shows typical weld settings for various SW stud diameters. The ideal settings in each weld position for a particular stud and workpiece should be established by beginning with typical settings and then varying them within the allowable range to meet the required conditions. Typical settings are based on good weld conditions, such as a clean workpiece, good ground, and others. Also necessary is a proper dc power source, which should have the following characteristics: • • • •

HIGH OPEN-CIRCUIT VOLTAGE (70 TO 100 V) RAPID OUTPUT CURRENT RISE DROOPING OUTPUT VOLT-AMPERE CURVE HIGH CURRENT OUTPUT OVER A SHORT TIME

Stud arc welding requires a very short weld time (Table 3), but a very high current input, when compared with other types of arc welding. Usually, the duty cycle for this process is much lower than that of other welding processes. Constantvoltage types of dc power sources are not suitable for stud welding. Many of the available dc power sources are acceptable for stud welding, including those supplied as integrated power/control systems by the stud-welding manufacturer. These integrated systems use both three-phase and single-phase incoming alternating current (ac) power. Both standard and special voltage units are available. Single-phase units are relatively low cost, portable systems that are usually suitable for smaller-diameter ( ≤ 12.7 mm, or

1 in.) stud arc welding. Three-phase units are preferred for larger2

diameter studs, because they provide a balanced load on incoming power lines and have smoother arc characteristics. A transformer/rectifier power source, combined with an integral stud-welding control system, is shown in Fig. 2. Weld cable length and cable size also can affect the total output of a power source. Basically, a larger welding cable size minimizes current loss that is due to cable length, resistance, and heating. Thus, when the distance from the power source to the gun increases significantly, or when the number and size of studs applied per minute are large, larger welding cables should be used. Smaller dc power sources suitable for stud arc welding can be wired in parallel to produce the necessary current range. Ferrules, or ceramic arc shields, are used in most stud arc weld applications. They are available in a wide variety of sizes

and shapes to fit specific stud base designs and applications. For example, there are specific designs for welding studs through metal deck, to the fillet or heel of an angle, to round pipes or bars, or in a vertical position. A ferrule is placed over every stud at the weld end, where it is held in place by a grip on the stud-welding gun. The ferrule is used only once and must be removed from the stud after welding is completed to allow the inspection of the completed weld. Typical ferrule configurations are shown in Fig. 8.

FIG. 8 CROSS SECTIONS OF FERRULES SHOWING SOME OF THE MANY VARIETIES AVAILABLE FOR DIFFERENT STUD BASE GEOMETRIES, WELD POSITIONS, AND APPLICATIONS

Ferrule design is important, because it controls certain functions during welding. For example: •

• • •

VENTS ON THE BOTTOM OF THE FERRULE ALLOW WELD GASES TO ESCAPE FROM THE WELD AREA AND RESTRICT THE INFLOW OF AIR TO MINIMIZE WELD POROSITY AND OXIDATION. THE INTERNAL CAVITY OF THE FERRULE CONFINES AND SHAPES THE MOLTEN METAL INTO A FLASH (FILLET) AROUND THE STUD PERIPHERY. HEAT IN THE WELD AREA IS CONCENTRATED AND CONTAINED. FLASH AND WELD SPATTER ARE MINIMIZED.

Although the dimensions of the flash are controlled by the ferrule configuration, the flash diameter and height must be taken into consideration when designing mating parts. Manufacturer specifications on finished stud flash dimensions should be followed, and test welds should be made to establish part fit. Flash dimensions can be accommodated by any of the five methods shown in Fig. 9.

FIG. 9 COMMON METHODS FOR ACCOMMODATING WELD FLASH (FILLET) WITH STUD ARC WELDS. TYPICAL DIMENSIONS FOR A, B, AND C ARE GIVEN IN TABLE 4.

The typical dimensions for flash clearance are shown in Table 4. It should be noted that the flash (fillet) of a stud weld is not the same as the fillet that is produced by conventional welding techniques. It may have areas of nonfusion on its vertical leg or shrink fissures, which have no adverse affect on weld strength or ductility (Ref 4). Consequently, it is not subject to the profile and inspection criteria used for a conventional weld fillet.

TABLE 4 DIMENSIONS FOR COUNTERBORE AND COUNTERSINK WELD FLASH CLEARANCE A, B, and C are shown in Fig. 9.

STUD BASE DIAMETER COUNTERBORE A B mm in. mm in. mm in. 1 6.4 11.1 0.437 3.2 0.125 7.9 9.5 11.1 12.7 15.9

4 5 16 3 8 7 16 1 2 5 8

90° COUNTERSINK C mm in. 3.2 0.125

12.7 0.500 3.2

0.125 3.2

0.125

15.1 0.593 3.2

0.125 3.2

0.125

16.7 0.656 4.7

0.187 3.2

0.125

19.1 0.750 4.7

0.187 4.7

0.187

22.2 0.875 5.5

0.218 4.7

0.187

19.1

3 4

28.6 1.125 7.9

0.312 4.7

0.187

Note: Dimensions can vary, depending on stud style and ceramic arc shield selected. Consult manufacturer for details. Base Plate Material and Thickness. To obtain full strength, the stud arc welding of low-carbon steel, stainless steel, or aluminum studs should be made to a base material of sufficient thickness so that the stud fails in tension or torque, rather than by pulling a hole in the base-metal workpiece. In the case of stainless or low-carbon steel studs, the base-metal thickness should be at least one-third that of the stud diameter. For aluminum studs, the aluminum base plate material should be at least half that of the stud diameter. Thinner base plates can be used if strength is not the primary design characteristic that is required. On thinner materials, the studs will pull a hole in the plate at strength levels below their maximum. If the plate is too thin, then the stud will bum a hole through it. The recommended base plate thickness for fullstrength development and the minimum thicknesses required to prevent burning through the base plate are given in Table 5.

TABLE 5 RECOMMENDED BASE METAL THICKNESSES FOR STUD ARC WELDING STUD BASE DIAMETER

STEEL BASE METAL FOR FULL STRENGTH(A)

MINIMUM(B)

mm 4.8

mm 1.59

in. 0.062

mm in. mm 0.91 0.036 3.18

in. 0.125

mm 3.18

in. 0.125

3.18

0.125

1.21 0.048 3.18

0.125

3.18

0.125

3.18

0.125

1.52 0.060 4.76

0.187

3.18

0.125

4.76

0.187

1.90 0.075 4.76

0.187

4.76

0.187

4.76

0.187

2.28 0.090 6.35

0.250

4.76

0.187

4.76

0.187

3.04 0.120 6.35

0.250

6.35

0.250

6.35

0.250

3.68 0.145 . . .

...

...

...

7.94

0.318

4.70 0.185 . . .

...

...

...

9.53

0.375

6.35 0.250 . . .

...

...

...

9.53

0.375

9.53 0.375 . . .

...

...

...

6.4 7.9 9.5 11.1 12.7 15.9 19.1 22.2 25.4

in. 3 16 1 4 5 16 3 8 7 16 1 2 5 8 3 4 7 8

1

ALUMINUM BASE METAL WITHOUT BACKUP MINIMUM WITH BACKUP(C)

(A) NEAREST COMMONLY AVAILABLE BASE METAL THICKNESS. (B) STEEL WITHOUT BACKUP--WILL NOT DEVELOP FULL STUD STRENGTH. (C) METAL BACKUP REQUIRED TO PREVENT MELT-THROUGH OF BASE METAL The most widely used combinations of base plate and stud materials are shown in Table 6. With low-carbon steel studs, no preheat or postheat treatments are needed when welding to low-carbon steel or austenitic stainless steel base plate materials. As carbon content in the base plate increases into the medium range, heat treatment for stud welding may be required. Welding to high-carbon base plate is not suggested. When base plate/stud combinations are questionable, an application qualification test is suggested (Ref 5). Details on this procedure are provided in the section "Stud-Welding Quality Control, Qualification, and Inspection" in this article. For best results when welding stainless steel studs to lowcarbon steel plate (especially when working stress levels are very high or there are repetitive loading cycles), the studs should either be manufactured from annealed-in-process materials or postmanufacture annealed before use.

TABLE 6 TYPICAL COMBINATIONS OF BASE AND STUD METALS FOR STUD ARC WELDING

BASE METAL LOW-CARBON STEEL, AISI 1006 TO 1022(A) STAINLESS STEEL, 300 SERIES(B), 405, 410, AND 430 ALUMINUM ALLOYS, 5000 SERIES(C)

STUD METAL LOW-CARBON STEEL, AISI 1006 TO 1020; STAINLESS STEEL, 300 SERIES LOW-CARBON STEEL, AISI 1006 TO 1020; STAINLESS STEEL, 300 SERIES ALUMINUM ALLOYS, 5000 SERIES(C)

Source: Ref 1

(A) REFER TO ANSI/AWS D1.1-92, TABLE 4.1 (GROUPS I AND II) FOR APPROVED STEELS. (B) EXCEPT FOR THE FREE-MACHINING TYPE 303 STAINLESS STEEL. (C) REFER TO ANSI/AWS D1.2-89, TABLES 7.1 AND 7.4. Good welding practice requires that the base plate material at the weld spot be cleaned. Common contaminants that can result in unsatisfactory welds if not removed include paint, heavy mill scale, heavy rust, oxidation, oil or grease, and plating, such as galvanizing and anodizing on aluminum. Removal methods vary according to the contaminant, the base material, and the end use. Wire brushing and grinding may be satisfactory on heavy structural steel but can be destructive to thin-gage metals or anodized aluminum, because they may reduce the thickness to a point where bum through or lessthan-full-strength welds result. Suggested cleaning methods include grinding, wire brushing, or needle scaling on heavy materials; using solvents to clean grease and oil; and using a noncontaminating stainless wire brush or a nonmetallic foam disk to clean aluminum and milling. Because anodizing is difficult to remove from aluminum without reducing the metal thickness, consideration should be given to anodizing the base plate and stud after the welding operation. Similarly, the weld end of the stud should be free from materials that would contaminate the weld, including paint, rust, galvanizing, and others. Studs cannot be coated with any nonconductive material that would interfere with the flow of welding current. Light copper flashing and nickel or chrome plating usually do not cause welding problems. Because a high-amperage, short-time current is typical in the stud arc welding process, it is also important that the grounding spot(s) on the base material be clean, that the ground be tight, and that all cables be in good condition and have tight connections. The position of the ground(s) also can influence stud-welding quality. This depends on the geometry of the base-metal shape and is the result of electromagnetic arc deflection, or arc blow (Ref 6). Occurrences are more common when studs are very near a free edge or are on long and narrow base plates, hollow pipe sections, or irregular peripheries, and can be corrected by moving the ground to a central position or using multiple grounds. Because current flow is usually away from the ground toward the heavier or larger area of the base plate, arc blow is characterized by a lack of flash or fillet on the side of the stud nearer the edge, end, or smaller area. Stud Strength. Fasteners will develop full material strength when they have been stud arc welded to compatible base plate alloys. Tables 7, 8, and 9 show the tension and torque loads for various threaded stud diameters, based on minimum specified stud strengths (Ref 2, 3, and 6). For unthreaded stud fasteners, the yield strength and tensile or ultimate strength is calculated by:

YIELD STRENGTH = ASFY TENSILE STRENGTH = ASFU where As is the area of the stud shank (in.2), Fy is the specified stud material yield strength (minimum psi) and Fu is the specified stud material ultimate strength (minimum psi).

TABLE 7 MECHANICAL PROPERTIES OF LOW-CARBON STEEL STUD ARC WELDED FASTENERS Fasteners have 380 MPa (55 ksi) minimum ultimate strength and 345 MPa (50 ksi) minimum yield strength. STUD THREAD DIAMETER(A)

MEAN EFFECTIVE THREAD AREA(B)

YIELD TENSILE LOAD(C)

ULTIMATE TENSILE LOAD

YIELD TORQUE(D)

mm2 in.2

kn

lbf

kn

lbf

j

ULTIMATE TORQUE

ULTIMATE SHEAR LOAD(E)

10-24 UNC 10-32 UNF

11

0.017 3.8

850

4.2

935

3.6

ft · j lbf 32(F) 4.0

ft · lbf kn

lbf

35.1(F) 3.1

701

13

0.020 4.4

1,000

4.9

1,100

4.3

38(F)

4.7

41.3(F) 3.7

825

1 -20 4

21

0.032 7.1

1,600

7.8

1,760

9.1

6.7

9.9

7.3

5.9

1,320

23

0.036 8.0

1,800

8.8

1,980

10.2

7.5

11.3

8.3

6.6

1,485

34

0.052 11.6

2,600

12.7

2,860

18.4

13.6

20.2

14.9

9.5

2,145

37

0.058 12.9

2,900

14.2

3,190

20.5

15.1

22.5

16.6

10.6

2,393

50

0.078 17.3

3,900

19.1

4,290

33.1

24.4

36.3

26.8

14.3

3,218

57

0.088 19.6

4,400

21.5

4,840

37.3

27.5

41.1

30.3

16.1

3,630

68

0.106 23.6

5,300

25.9

5,830

52.3

38.6

57.6

42.5

19.5

4,373

76

0.118 26.2

5,900

28.9

6,490

58.3

43.0

64.1

47.3

21.7

4,868

92

0.142 31.6

7,100

34.7

7,810

80.3

59.2

88.3

65.1

26.0

5,856

103

0.160 35.6

8,000

39.1

8,800

90.4

66.7

99.4

73.3

29.4

6,600

146

0.226 50.3

11,300 55.3

41.5

9,323

UNC 1 -28 4

UNF 5 -18 16

UNC 5 -24 16

UNF 3 -16 8

UNC 3 -24 8

UNF 7 -14 16

UNC 7 -20 16

UNF 1 -13 2

UNC 1 -20 2

UNF 5 -11 8

12,430 159.6 117.7 175.5 129.5

UNC 5 -18 8

165

0.255 56.7

12,750 62.4

14,025 180.1 132.8 198.1 146.1

46.8

10,519

215

0.334 74.3

16,700 81.7

18,370 283.1 208.8 311.3 229.6

61.3

13,778

240

0.372 82.7

18,600 91.0

20,460 315.2 232.5 346.8 255.8

68.3

15,345

298

0.462 102.8 23,100 112.8 25,355 456.8 336.9 501.4 369.8

84.6

19,017

328

0.509 113.2 25,450 124.5 27,995 503.1 371.1 553.6 408.3

93.4

20,996

91 437

0.606 134.8 30,300 148.0 33,275 684.7 505.0 751.9 554.6 0.678 150.8 33,900 165.9 37,290 766.0 565.0 842.6 621.5

111.0 24,956 124.4 27,967

UNF 3 -10 4

UNC 3 -16 4

UNF 7 -9 UNC 8 7 -14 8

UNF 1-8 UNC 1-14 UNF

(A) UNC, UNIFIED COARSE THREAD SERIES; UNF, UNIFIED FINE THREAD SERIES. (B) MEAN EFFECTIVE THREAD AREA IS BASED ON A MEAN FULL DIAMETER MIDWAY BETWEEN MINOR AND PITCH THREAD DIAMETERS. (C) IN PRACTICE, A STUD SHOULD NOT BE USED AT OR HIGHER THAN YIELD LOAD. A FACTOR OF SAFETY SHOULD BE APPLIED, AND 60% OF YIELD IS COMMONLY USED, ALTHOUGH OTHER VALUES MAY BE USED AT DISCRETION OF USER. (D) TORQUE FIGURES ARE BASED ON THE ASSUMPTION THAT EXCESSIVE THREAD DEFORMATION HAS NOT AFFECTED THE PROPORTIONAL RANGE OF TORQUE/TENSION RELATIONSHIP. AN AVERAGE TORQUE COEFFICIENT OF 0.20 WAS USED IN THESE CALCULATIONS. (E) SHEAR LOAD IS BASED UPON 0.75 TIMES ULTIMATE LOAD. THE USER SHOULD APPLY AN APPROPRIATE SAFETY FACTOR TO THESE FIGURES. (F) VALUE GIVEN IN IN. · LBF. TABLE 8 MECHANICAL PROPERTIES OF STAINLESS STEEL STUD ARC WELDED FASTENERS Fasteners have 520 MPa (75 ksi) minimum ultimate strength and 205 MPa (30 ksi) minimum yield strength. STUD THREAD DIAMETER(A)




YIELD TORQUE(D)

mm2

kn

kn

j

in.2

lbf

lbf

ULTIMATE TORQUE

ft · lbf (F)

j

5.4


ft · lbf (F)

47.8

kn

lbf

4.3

956

10-24 UNC 10-32 UNF

11

0.017 2.3

510

5.7

1,275

2.2

19.1

13

0.020 2.7

600

6.7

1,500

2.5

22.5(F) 6.4

56.3(F) 5.0

1,125

1 -20 4

21

0.032 4.3

960

10.7

2,400

5.4

4.0

13.6

10.0

8.0

1,800

23

0.036 4.8

1,080

12.0

2,700

6.1

4.5

15.3

11.3

9.0

2,025

34

0.052 6.9

1,560

17.3

3,900

11.0

8.1

27.5

20.3

13.0

2,925

UNC 1 -28 4

UNF 5 -18 16

UNC

5 -24 16

37

0.058 7.7

1,740

19.3

4,350

12.3

9.1

30.6

22.6

14.5

3,263

50

0.078 10.4 2,340

26.0

5,850

19.8

14.6

49.6

36.6

19.5

4,388

57

0.088 11.7 2,640

29.4

6,600

22.4

16.5

56.0

41.3

22.0

4,950

68

0.106 14.1 3,180

35.4

7,950

31.5

23.2

78.6

58.0

26.5

5,963

76

0.118 15.7 3,540

39.4

8,850

35.0

25.8

87.4

64.5

29.5

6,638

92

0.142 18.9 4,260

47.4

10,650 48.1

35.5

120.4

88.8

35.5

7,988

103

0.160 21.4 4,800

53.4

12,000 54.2

40.0

135.6

100.0

40.0

9,000

146

0.226 30.2 6,780

75.4

16,950 95.7

70.6

239.4

176.6

56.5

12,713

165

0.255 34.0 7,650

85.1

19,125 108.1 79.7

270.1

199.2

63.8

14,344

215

0.334 44.6 10,020 111.4 25,050 169.9 125.3

424.50 313.1

83.6

18,788

240

0.372 49.6 11,160 124.1 27,900 189.1 139.5

472.9

348.8

93.1

20,925

298

0.462 61.7 13,860 154.1 34,650 274.0 202.1

685.1

505.3

115.6 25,988

328

0.509 67.9 15,270 169.8 38,175 301.9 222.7

754.8

556.7

127.4 28,631

391 437

0.606 80.9 18,180 202.2 45,450 410.8 303.0 0.678 90.5 20,340 226.2 50,850 459.6 339.0

1027.0 757.5 1149.0 847.5

151.6 34,088 169.6 38,138

UNF 3 -16 8

UNC 3 -24 8

UNF 7 -14 16

UNC 7 -20 16

UNF 1 -13 2

UNC 1 -20 2

UNF 5 -11 8

UNC 5 -18 8

UNF 3 -10 4

UNC 3 -16 4

UNF 7 -9 UNC 8 7 -14 8

UNF 1-8 UNC 1-14 UNF

(A) UNC, UNIFIED COARSE THREAD SERIES; UNF, UNIFIED FINE THREAD SERIES. (B) META, MEAN EFFECTIVE THREAD AREA, IS BASED ON A MEAN FULL DIAMETER MIDWAY BETWEEN MINOR AND PITCH THREAD DIAMETERS. (C) IN PRACTICE, A STUD SHOULD NOT BE USED AT OR HIGHER THAN YIELD LOAD. A FACTOR OF SAFETY SHOULD BE APPLIED, AND 60% OF YIELD IS COMMONLY USED, ALTHOUGH OTHER VALUES MAY BE USED AT THE DISCRETION OF USER. (D) TORQUE FIGURES ARE BASED ON THE ASSUMPTION THAT EXCESSIVE THREAD DEFORMATION HAS NOT AFFECTED THE PROPORTIONAL RANGE OF TORQUE/TENSION RELATIONSHIP. AN AVERAGE TORQUE COEFFICIENT OF 0.20 WAS USED IN THESE CALCULATIONS. (E) SHEAR LOAD IS BASED UPON 0.75 TIMES ULTIMATE LOAD. THE USER SHOULD APPLY AN APPROPRIATE SAFETY FACTOR TO THESE FIGURES.

(F) VALUE GIVEN IN IN. · LBF. TABLE 9 MECHANICAL PROPERTIES OF ALUMINUM STUD ARC WELDED FASTENERS Fasteners have 290 MPa (42 ksi) minimum ultimate strength and 205 MPa (30 ksi) minimum yield strength. STUD THREAD DIAMETER(A)

10-24 UNC 10-32 UNF 1 -20 UNC 4 1 -28 UNF 4 5 -18 UNC 16 5 -24 UNF 16 3 -16 UNC 8 3 -24 UNF 8 7 -14 UNC 16 7 -20 UNF 16 1 -13 UNC 2 1 -20 UNF 2




YIELD TORQUE(D)

ULTIMATE TORQUE


mm2 11

in.2 0.017

kn 2.3

lbf 510

kn 3.2

lbf 714

j 2.2

13 21

0.020 0.032

2.7 4.3

600 960

3.9 6.0

840 1344

23

0.036

4.8

1080

6.7

34

0.052

6.9

1560

37

0.058

7.7

50

0.078

57

j 3.0

ft · lbf 26.8(F)

kn 1.9

lbf 428

2.5 5.4

ft · lbf 19. 1(F) 22.5(F) 4.0

3.6 7.6

31.5(F) 5.6

2.2 3.6

504 806

1512

6.1

4.5

8.5

6.3

4.0

907

9.7

2184

11.0 8.1

15.5

11.4

5.8

1310

1740

10.8

2436

12.3 9.1

17.2

12.7

6.5

1462

10.4

2340

14.6

3276

19.8 14.6

27.8

20.5

8.7

1966

0.088

11.7

2640

16.4

3696

22.3 16.5

31.3

23.1

9.6

2218

68

0.106

14.1

3180

19.8

4452

31.5 23.2

44.1

32.5

11.9

2671

76

0.118

15.7

3540

22.0

4956

35.0 25.8

48.9

36.1

13.2

2974

92

0.142

18.9

4260

26.5

5964

48.1 35.5

67.4

49.7

15.9

3578

103

0.160

21.4

4800

29.9

6720

54.2 40.0

75.9

56.0

17.9

4032

(A) UNC, UNIFIED COARSE THREAD SERIES; UNF, UNIFIED FINE THREAD SERIES. (B) MEAN EFFECTIVE THREAD AREA IS BASED ON A MEAN FULL DIAMETER MIDWAY BETWEEN MINOR AND PITCH THREAD DIAMETERS. (C) IN PRACTICE, A STUD SHOULD NOT BE USED AT OR HIGHER THAN YIELD LOAD. A FACTOR OF SAFETY SHOULD BE APPLIED, AND 60% OF YIELD IS COMMONLY USED, ALTHOUGH OTHER VALUES MAY BE USED AT DISCRETION OF USER. (D) TORQUE FIGURES ARE BASED ON THE ASSUMPTION THAT EXCESSIVE THREAD DEFORMATION HAS NOT AFFECTED THE PROPORTIONAL RANGE OF TORQUE/TENSION RELATIONSHIP. AN AVERAGE TORQUE COEFFICIENT OF 0.20 WAS USED IN THESE CALCULATIONS. (E) SHEAR LOAD IS BASED ON 0.60 TIMES ULTIMATE LOAD. THE USER SHOULD APPLY AN APPROPRIATE SAFETY FACTOR TO THESE FIGURES. (F) VALUE GIVEN IN IN. · LBF. For threaded studs, the area is based on the mean effective thread area, which is calculated by:

AS = 0.7854 [D- (0.9743/N)]2

where D is the nominal diameter of the stud and N is the number of threads per inch. Process Variations and Special Equipment. There are several specific applications that lend themselves to special

variations of the stud arc welding technique. One application is the welding of studs to thin base materials that are less than the minimum thicknesses listed in Table 5. Although full weld base strength is usually not achieved, the resulting strength level is suitable for the application loadings involved. This process variation, which is called short-cycle stud arc welding, does not use a ceramic ferrule. Instead, higher weld current is used with very short times, which minimizes penetration of the stud into the base plate. Normally, this variation is used with stud diameters ≤ 9.5 mm ( ≤ 0.375 in.) in situations where backside marking is not considered detrimental. The second special process is gas-shielded arc welding, which also does not use a ceramic ferrule. Instead, the studwelding area is shielded by an inert gas, usually argon. This process can be used with both steel and aluminum, but is more widely used with the latter material. The welding variables fall into a very narrow range, and conditions for application usually include tight tolerances and a tightly controlled setup. Consequently, the stud arc welding equipment most often consists of a fixed-gun production unit. A typical application is the welding of aluminum studs with special end configurations to aluminum kitchen utensils. Special equipment is frequently used with the short-cycle and/or gas-shielded arc welding processes, because the application parameters must be tightly controlled and involve large quantities of studs applied in a production environment. Equipment manufacturers assemble many types of special equipment for production-line use. The equipment can be used with automatic stud feed, automatic stud and ferrule feed, computer-programmed indexing, robotic stud-welding guns, and other mechanisms. A typical automated stud arc welding production unit is shown in Fig. 10. A unit can involve simple, column-mounted, single-gun systems or sophisticated, multigun, multiple feed units that cost thousands of dollars.

FIG. 10 TYPICAL AUTOMATIC STUD-FEED HAND-GUN SYSTEM FOR PORTABLE OPERATION USING TEMPLATES OR FIXTURES ON THE WORKPIECE


1. "RECOMMENDED PRACTICES FOR STUD WELDING," ANSI/AWS C5.4, AWS 2. "STRUCTURAL WELDING CODE--STEEL," ANSI/AWS D1.1, SECTION 7, "STUD WELDING," AWS 3. "STRUCTURAL WELDING CODE--ALUMINUM," ANSI/AWS D1.2, AWS 4. "STRUCTURAL WELDING CODE--STEEL," ANSI/AWS D1.1, SECTION 7, "STUD WELDING," FOOTNOTE 27, AWS 5. "STRUCTURAL WELDING CODE--STEEL," ANSI/AWS D1.1, SECTION 7, "STUD WELDING," PARAGRAPH 7.6, AWS 6. W.A. BAESLACK, G. FAYER, S. REAM, AND C.E. JACKSON, QUALITY CONTROL IN ARC STUD WELDING, WELD. J., NOV 1975, P 789-798 Stud Arc Welding Harry A. Chambers, TRW Nelson Stud Welding Division

Fixturing and Tooling for Stud Arc Welding Regardless of the stud-welding method employed, certain "expendable" accessories are required. These include such items as a chuck or collet to hold the particular stud being welded, a ferrule grip for the fastener, a foot that holds the ferrule grip, legs that attach to the gun and adjust to accommodate various stud lengths, and other minor items. The operating life of these items is variable, depending on the care and maintenance they receive. For example, a chuck may last for 5000 to 25,000 welds. Legs last indefinitely, whereas ferrule grips and feet usually have a shorter usage life. Accessories are relatively inexpensive items that should be planned for when preparing budgetary estimates on any given project. The extent and sophistication of tooling for stud welding reflects the required production rate and the total number of studs to be welded. Locating the stud arc welding centers can be as simple as laying out the workpiece and center punching the locations either directly or through a template, as shown in Fig. 11. The studs are then placed in the punch marks, the stud-welding gun is held vertically, and the weld is initiated. Although operator skill is a factor, careful welders can achieve a perpendicularity of ±5° and a location tolerance of approximately 1.2 mm (0.046 in.). Often, the cover plate in a base plate/cover plate assembly can be used as the template for marking.

FIG. 11 CENTER PUNCHING BASE MATERIAL THROUGH A TEMPLATE

A permanent template becomes more practical when an increasing number of studs are to be welded on repetitive locations. Stud arc welded fasteners can be welded directly through temperature-resistant material which should be spaced off the work by 1.65 to 6.35 mm (0.065 to 0.250 in.) to allow expelled weld gases and weld spatter to escape without restriction, which could adversely affect weld quality. This type of template arrangement is shown in Fig. 12. Note that the template holes are drilled slightly larger than the outside dimension of the ceramic ferrule. Before the

template is prepared, it is good practice to consult the manufacturer specifications for the ceramic ferrule dimensions to be used with the stud being welded. Location accuracy, in this case, can be as tight as ±0.78 mm (±0.031 in.).

FIG. 12 WELDING THROUGH A TEMPLATE USING A FERRULE WITH STUD ARC WELDING FASTENERS

Even greater accuracy can be obtained, both in the vertical and horizontal directions, by adapting the template hole to fit a drill jig bushing tightly and by inserting the template adapter (or a ferrule tube adapter) through the bushing to the workpiece. The ferrule tube adapter also can be used for welding down into deeply drilled holes or into tight-fit areas (Fig. 13).

FIG. 13 PORTABLE STUD ARC WELDING GUN EQUIPPED WITH ACCESSORIES FOR WELDING STUDS THROUGH

OR INTO A DRILLED OR FORMED HOLE

With closely spaced studs, the foot grip is sometimes modified so that after welding the initial stud at a center-punched location, the foot can be held against the reference stud to weld the next or adjacent stud on the required spacing. The verticality of the weld stud can be ensured by several methods. The most accurate is obviously a fixed gun on a slide mechanism of the type used for automatic-feed, high-production work. This type is also the most expensive. A portable gun can be mounted on a machine slide or drill press by fabricating a bracket (Fig. 14) in cases where production quantities do not justify the expense of automatic equipment.

FIG. 14 INEXPENSIVE METHOD OF MOUNTING PORTABLE HAND GUN TO MACHINE SLIDE OR DRILL PRESS TO OBTAIN FIXED-POSITION ACCURACY

Other methods of ensuring verticality are to use a bushing-type template, a bubble level mounted on the rear of the gun, or a bipod foot arrangement, as shown in Fig. 15. With the stud and ferrule acting as one fixed point, the two bipod screws are adjusted to obtain perpendicularity. Vertical variations that are less than ±1° can be achieved with proper accessories and adjustments.

FIG. 15 USE OF AN ADJUSTABLE BIPOD FOOT ASSEMBLY TO PROVIDE THREE-POINT CONTACT, ENSURING PERPENDICULARITY OF STUD TO WORKPIECE

A wide range of accessories is available for welding studs of different lengths or for welding studs into areas with limited access. Two widely used special accessories are shown in Fig. 16 and 17. For unusual situations, the manufacturer should be consulted to suggest or, possibly, design stud-welding accessories.

FIG. 16 TYPICAL OFFSET ASSEMBLY USED WHEN STUD ARC WELDED FASTENERS MUST BE WELDED BELOW AN OBSTRUCTION

FIG. 17 TYPICAL LEG-EXTENSION ACCESSORY FOR WELDING LONG STUDS (UP TO 1.83 M, OR 6 FT). NOTE THE USE OF A SPLIT FOOT AND FERRULE GRIP; THE GUN DOES NOT HAVE TO BE STRIPPED THE FULL LENGTH OF THE STUD WHEN THE STUD IS WELDED.


Stud-Welding Quality Control, Qualification, and Inspection Because of its broad use for many years and the relative simplicity of its applications, stud welding in the downhand position is considered to be a prequalified procedure, that is, it is only necessary to weld two studs satisfactorily in the downhand position to qualify both the process and the operator. The two qualification studs can be bend tested, torque tested, or tension tested. The physical test is satisfactory if no failure occurs in the weld or the heat-affected zone. Stud-welding code requirements (for example, Ref 2 and 7) define the necessary tests for prequalified stud welding to the flat or down-hand position, as well as other positions. The limit on the flat position is defined as 0 to 15° slope on the

surface to which the stud is applied. Beyond 15°, a ten-stud test is required. Again, each of the ten studs in the nonflat position must be bend, tensile, or torque tested. Bend testing is 90° from the original axis, whereas torque and tension testing are conducted until failure. Failure must occur in the stud shank or in the workpiece material, not in the weld or the heat-affected zone. All failures require retesting. Satisfactory testing qualifies the welder and the process. Typical bend, torque, and tensile testing setups are shown in Fig. 18, 19, and 20. Qualification records, including stud drawings, ceramic arc shield drawings (if used), weld settings, and position should be kept.

FIG. 18 TYPICAL BEND TESTING SETUP. STUDS THAT MUST UNDERGO APPLICATION QUALIFICATION SHOULD BE BENT 90° WITHOUT WELD FAILURE (15° FOR ALUMINUM STUDS). TESTING FOR PREPRODUCTION OR TESTING DURING PRODUCTION FOR INSPECTION PURPOSES REQUIRES BEND TESTING TO 30° FROM ORIGINAL POSITION (15° FOR ALUMINUM STUDS). FOR THREADED STUDS, TORQUE TESTING PER FIG. 19 SHOULD BE USED. SOURCE: REF 2

FIG. 19 TYPICAL TORQUE TESTING ARRANGEMENT. STUDS THAT MUST UNDERGO APPLICATION QUALIFICATION OR PREPRODUCTION TESTING SHOULD BE TORQUED TO DESTRUCTION WITHOUT FAILURE IN THE WELD. PROOF TORQUE TESTING DURING INSPECTION OF THREADED STUDS SHOULD BE APPLIED AT 60 TO 66% OF THE YIELD TORQUE LOAD FOR THE STUD SIZE TESTING. REFER TO TABLES 7, 8, AND 9. SOURCE: REF 2

FIG. 20 TYPICAL TENSION TEST FIXTURE. FOR QUALIFICATION OF APPLICATION OR PREPRODUCTION TESTING, EITHER BEND OR TORQUE TESTING CAN BE SUBSTITUTED FOR TENSION TESTING.

Either two-stud or ten-stud testing should be conducted on a production piece or on a separate piece of the same type and thickness (±25%) of material that will be used in the application, in the same position in which the studs will actually be applied. In addition to physical testing, the studs that are welded for qualification also must be visually inspected, whether there are two or ten. Visual inspection consists of verifying that a full 360 ° weld flash (fillet) is formed around the stud base. There should be no gaps or undercuts in the flash or in the wetted area. As mentioned previously, flash can be irregular in both height and evenness and it can contain shrinkage fissures or cracks. Because flash is expelled, not deposited, metal, it should not be subject to the usual inspection criteria for a weld fillet. The final visual inspection should be a measurement of the after-weld length. With stud arc welding, the after-weld length of the stud will be shorter by the lengths shown in Table 3. A consistent after-weld length, combined with acceptable visual inspection and physical testing, are assurances of satisfactory welds. Figure 21 shows typical stud arc weld appearances for both acceptable and problematic welds.

FIG. 21 STUD ARC WELD CHARACTERISTICS FOR SATISFACTORY AND UNSATISFACTORY WELDS. (A) SATISFACTORY; GOOD FLASH FORMATION, WETTING ENTIRE PERIPHERY OF STUD BASE. (B) UNSATISFACTORY; PLUNGE TOO SHORT, NOT ENOUGH MATERIAL ALLOWED FOR BURN-OFF. STUD MUST PROJECT AN ADEQUATE LENGTH BEYOND THE FERRULE. THIS CONDITION CAN ALSO BE CAUSED BY BAD GROUNDING PRACTICE. (C) UNSATISFACTORY; HANG UP CAUSED BY INADEQUATE ARC LENGTH LIFT OR BY MISALIGNMENT OF STUD, RESTRICTING FREE MOVEMENT. (D) UNSATISFACTORY; MISALIGNMENT OF STUD BY POOR GUN POSITION. (E) UNSATISFACTORY; COLD WELD, RESULTING IN NONUNIFORM OR INCOMPLETE FLASH AROUND STUD PERIPHERY. (F) UNSATISFACTORY; HOT WELD, WHERE MOLTEN MATERIAL IS EXPELLED VIOLENTLY FROM WELD, RESULTING IN POOR FLASH FORMATION.

Passing the two-stud preproduction or the ten-stud application qualification test approves the procedure and operator. Production stud welding is then allowed to proceed. However, at the start of each new production period, a two-stud test is again required before production can begin. Similarly, a change in stud diameter, equipment, settings, or operator requires requalification.

During production, welding inspection and testing should be used on a continuous basis to confirm acceptable quality. This can be done on a reasonable number of studs, depending on the welding conditions and stud appearance. After the visual inspection of all stud welds, physical testing of at least one weld per hundred is usually suggested. Any weld that does not show a full 360° flash or wetness should be tested. Welds without a full flash can be repair welded, but repaired welds are also required to be physically tested. In the case of unthreaded studs, the quality inspection test can be to bend 15°. The studs can be straightened after testing, if required. Threaded studs can be subjected to a torque test to approximately 80% of the yield torque load, rather than a bend test. Failure in the weld or heat-affected zone is cause to reject the weld being tested. Failure also requires that the quantity of studs welded from the previous test to the present test be carefully inspected for weld deficiencies. Additional tests may be necessary to confirm weld quality. If the quality inspection test studs must be straightened for end use, then they should be returned to the vertical position by using a bending tool. Continuous slow pressure should be applied during the straightening process. Temperature at the time of stud welding and during stud testing is an influential parameter. Stud welding should not be conducted at temperatures below -18 °C (0 °F). When the temperature is below 0 °C (32 °F), one additional stud per hundred welded should be tested to verify weld quality. Below 10 °C (50 °F), welded studs should be tested using the bending tool, rather than by hammer testing. In colder welding temperature conditions, threaded studs should always be tested by torque, rather than by bending. Impact testing at low temperatures produces brittle failures. This is not a function of weld quality, but of the ambient temperature at the time of welding and testing. Data resulting from extensive tests conducted on studs welded and tested at various temperatures have shown that it is the temperature and manner of testing that influences the failure mode, rather than the temperature at the time of welding (Ref 8). Another factor, besides temperature, that can affect weld quality is cleanliness of the weld area. Although stud arc welding does provide some cleaning action during welding, it is not sufficient to remove heavy contaminants. Coatings or rust on studs also can cause poor welds. Plating, except for a copper flash, nickel plating, or chromium plating, should be removed from the stud end prior to welding. This includes such materials as paint, zinc, and cadmium. Stud arc welded fasteners are usually capped during plating so that the weld end is left clean. The plating or contaminant removal methods should be compatible with the base-material thickness and composition. For example, a heavy paint or rust cover on thick structural steel can be removed by grinding, needle scaling, sand blasting, or chipping, and the weld surface will still be acceptable for stud welding. Stainless steel, brass, aluminum, or other polished surfaces should be cleaned with appropriate solvents or other cleaning methods that will not destroy the material finish or contaminate the weld area. For example, a carbon-steel wire wheel or brush should not be used to clean stainless steel or aluminum base material. Studs and ferrules should be kept clean and dry prior to welding. Although cold conditions will not adversely affect either, they should be protected from exposure to adverse weather conditions, high humidity, marine atmospheres, and other extremes. When ferrules absorb water, the steam generated during the welding process can cause very poor welds and excessive weld spatter. Ferrules exposed to moisture should be oven dried at a temperature between 101 and 120 °C (215 and 250 °F) for 2 h before use. The condition of weld cables, connections, and ground clamps can affect weld quality. These items should be checked frequently and either repaired or replaced, so that the cables are not frayed, connections are tight and clean, and ground clamps are properly functioning. Ground clamps should be tightly connected to the workpiece or work platen and attached to areas that have been cleaned of all deleterious materials that would prevent good current flow.


2. "STRUCTURAL WELDING CODE--STEEL," ANSI/AWS D1.1, SECTION 7, "STUD WELDING," AWS 7. "WELDED STEEL CONSTRUCTION (METAL ARC WELDING), STANDARD W59," CANADIAN STANDARDS ASSOCIATION 8. KENNEDY AND D.J. LAURIE, STUD WELDING AT LOW TEMPERATURES, CAN. J. CIVIL ENG., VOL 7, 1980, P 442-455


Stud-Welding Safety Precautions Any welding process can be dangerous if proper safety precautions are not followed. Equipment should be properly installed according to the directions of the manufacturer and maintained in good condition. Operators should be thoroughly trained in the use of the stud-welding process. They also should be familiar with the installation, operation, and maintenance procedures for the equipment being used. Electrical shock is a potential cause of injury or death. Stud-welding systems should be properly installed and grounded using electrical connections made according to national (Ref 9) and local code guidelines. Neither operators nor anyone else in the welding area should ever touch live electrical parts, or weld in wet areas, or wear wet gloves or clothing. In addition, all workers should ensure that they are insulated from shock. All electric cables and connections should be kept in good condition. They should be inspected on a regular basis and frayed sections, broken insulation, or broken connectors should be repaired or replaced at once. Stud welding frequently uses long cable lengths, particularly on construction sites, where the cables are subject to abuse by towmotors, trucks, material storage, or foot traffic. It should be ensured that these conditions do not result in cable damage. Although weld spatter and arc flash are minimal with the stud-welding process, they do occur and precautions should be taken. All combustible or volatile materials should be removed from the weld area so that sparks or spatter cannot reach them. Gas cylinders should be stored and secured properly and checked to ensure that they do not contact any electrical cables in the welding circuit. Proper protective clothing should be worn, including boots, aprons, and gloves, as necessary. Eye protection is always necessary. Eye glasses with spectacle frames, side shields, and lenses with shade number three absorption filters should be worn by the operator at all times (Ref 10). Helpers or workers within 1.5 m (5 ft) of the weld area also should wear clear safety glasses with side shields. Fire-suppression equipment should be available in or adjacent to the weld area for immediate use in emergencies. Ventilation of the welding area is necessary. Fumes from welding and cleaning materials, such as solvents, as well as paints, epoxies, and galvanizing or other coatings, can be harmful. Ventilation can be either forced or natural, depending on the job conditions. Material suppliers should provide material safety data sheets on any items being used in the welding area, and proper precautions for potentially dangerous contents should be followed. Pinch points where fingers or hands can be caught should be avoided during material handling or where moving parts are involved in the welding process. The instructions of the manufacturer should be followed when maintaining and servicing the stud-welding equipment. If possible, all electric supplies should be disconnected and locked out during maintenance or troubleshooting work. Finally, certain welding operations produce elevated noise levels during the weld cycle. Operators and other workers near the equipment should use hearing protection that can adequately protect them, in accordance with Ref 11.


9. "NATIONAL ELECTRICAL CODE" NFPA-70, NATIONAL FIRE PROTECTION ASSOCIATION 10. "SAFE PRACTICE FOR OCCUPATION AND EDUCATIONAL EYE AND FACE PROTECTION," Z87.1, ANSI 11. "WELDING, CUTTING, AND BRAZING," STANDARD 29CFR, PART 1910, OCCUPATIONAL SAFETY AND HEALTH ADMINISTRATION Stud Arc Welding

Harry A. Chambers, TRW Nelson Stud Welding Division

References

1. "RECOMMENDED PRACTICES FOR STUD WELDING," ANSI/AWS C5.4, AWS 2. "STRUCTURAL WELDING CODE--STEEL," ANSI/AWS D1.1, SECTION 7, "STUD WELDING," AWS 3. "STRUCTURAL WELDING CODE--ALUMINUM," ANSI/AWS D1.2, AWS 4. "STRUCTURAL WELDING CODE--STEEL," ANSI/AWS D1.1, SECTION 7, "STUD WELDING," FOOTNOTE 27, AWS 5. "STRUCTURAL WELDING CODE--STEEL," ANSI/AWS D1.1, SECTION 7, "STUD WELDING," PARAGRAPH 7.6, AWS 6. W.A. BAESLACK, G. FAYER, S. REAM, AND C.E. JACKSON, QUALITY CONTROL IN ARC STUD WELDING, WELD. J., NOV 1975, P 789-798 7. "WELDED STEEL CONSTRUCTION (METAL ARC WELDING), STANDARD W59," CANADIAN STANDARDS ASSOCIATION 8. KENNEDY AND D.J. LAURIE, STUD WELDING AT LOW TEMPERATURES, CAN. J. CIVIL ENG., VOL 7, 1980, P 442-455 9. "NATIONAL ELECTRICAL CODE" NFPA-70, NATIONAL FIRE PROTECTION ASSOCIATION 10. "SAFE PRACTICE FOR OCCUPATION AND EDUCATIONAL EYE AND FACE PROTECTION," Z87.1, ANSI 11. "WELDING, CUTTING, AND BRAZING," STANDARD 29CFR, PART 1910, OCCUPATIONAL SAFETY AND HEALTH ADMINISTRATION Stud Arc Welding Harry A. Chambers, TRW Nelson Stud Welding Division

Selected References

• "DESIGN AND APPLICATION HAND BOOK FOR LIGHT GAUGE METALWORKING," PUBLICATION CD-92, TRW NELSON STUD WELDING DIVISION, 1992 • "INDUSTRIAL DESIGN DATA--NELSON STUD WELDING PROCESS," PUBLICATION MB-12-86, TRW NELSON STUD WELDING DIVISION • "SAFETY IN WELDING AND CUTTING," ANSI/AWS Z49.1, AWS • T.E. SHOUP, "STUD WELDING," BULLETIN 214, WELDING RESEARCH COUNCIL, APRIL 1976 • "TOOLING TECHNIQUES FOR STUD WELDING," TRW NELSON STUD WELDING DIVISION • STUD WELDING, CHAPTER 9, WELDING HANDBOOK, 8TH ED., AWS, 1990

Capacitor Discharge Stud Welding Richard L. Alley, American Welding Society

Introduction CAPACITOR DISCHARGE STUD WELDING is a stud arc welding process in which the tip of the stud melts almost instantly when energy stored in capacitors is discharged through it. The three basic modes of capacitor discharge (CD) stud welding are initial-gap welding, initial-contact welding, and drawn-arc welding. Figure 1 shows current-versus-time curves for the three process variations.

FIG. 1 TYPICAL CURRENT-VERSUS-TIME CURVES FOR THE THREE CD STUD WELDING METHODS. SOURCE: REF 1

The initial-gap mode (Fig. 2) is begun with the stud held away from the work surface by the welding head. At the beginning of the weld cycle, the stud is forced against the work surface, melting the tip, the face of the stud, and the adjoining work surface upon contact with the work surface. The weld is completed using the gun forces (i.e., spring pressure or air pressure) to plunge the stud into the molten materials, forming a strong welded bond between the stud and the work surface. The weld cycle time (Fig. 1) for this process is from 4 to 6 ms, and the penetration of the weld zone into the work surface is normally from 0.10 to 0.15 mm (0.004 to 0.006 in.).

FIG. 2 INITIAL-GAP CD STUD WELDING. SEE TEXT FOR EXPLANATION. SOURCE: REF 1

The initial-contact mode (Fig. 3) begins with the weld stud in contact with the work surface. The weld cycle is initiated

with a surge of current that disintegrates the weld tip, thus melting the stud face area and the work surface that it immediately contacts. The stud is forced into the molten material, forming a strong homogeneous weld. This process has a weld cycle time of approximately 6 ms, much like the gap process.

FIG. 3 INITIAL-CONTACT CD STUD WELDING. SEE TEXT FOR EXPLANATION. SOURCE: REF 1

The drawn-arc mode (Fig. 4) begins with the stud in contact with the work surface. When the weld cycle is initiated, a current surge is applied to the weld tip and the stud is retracted from the work surface, drawing an arc that melts the face of the stud and the work surface directly beneath it. The stud is then plunged into the molten pool of material, forming a welded connection. The weld cycle time for this process is longer than for the other two processes, and the heat-affected zone (HAZ) is thicker than it is in the preceding two processes.

FIG. 4 DRAWN-ARC CD STUD WELDING. SEE TEXT FOR EXPLANATION. SOURCE: REF 1

Advantages and Disadvantages. A major reason for using the CD stud welding process is that it provides a strong welded fastener with either a minimum or no reverse side marking of the work material. Another reason is that it is cost effective, especially for small-diameter fasteners. This process also allows fasteners to be welded to very thin sections of work material, as well as to thick sections (as thick as necessary), with reliable results. Furthermore, the process allows the welding of many dissimilar material combinations, such as aluminum studs to zinc die castings. The disadvantages of using the CD process are the limited stud diameters available and the fact that the work surface must be clean of mill scale, dirt, oxidation products, and oil.

Reference

1. R.L. O'BRIEN, ED., WELDING HANDBOOK, 8TH ED., VOL 2, AWS, 1991, P 317-323 Capacitor Discharge Stud Welding Richard L. Alley, American Welding Society

Applications The CD process is normally used for fastener sizes of 4-40 through in.). Although there are limited uses of

1 1 -20, with lengths from 6.4 to 38.1 mm ( to 1.5 4 4

5 3 -18 and -16 diameter studs in special applications, only the more-common 16 8

stud diameters will be considered in this article. The work material thickness can typically range from 0.25 mm (0.010 in.) to as thick as the application requires. The determining factors for making successful welds are the stud diameter, the materials being used, the strength required, and the amount of allowable reverse side marking. See Table 1 for guidance on the expected strength parameters.

TABLE 1 STANDARD CAPACITOR DISCHARGE STUD FASTENER LOAD STRENGTHS STUD MATERIAL

STUD SIZE

LOW-CARBON COPPER-FLASHED STEEL

6-32 8-32 10-24 1 -20 4 5 -18 16 3 -16 8 6-32 8-32 10-24 1 -20 4 5 -18 16 3 -16 8 6-32 8-32 10-24 1 -20 4 5 -18 16 3 -16 8 1 -13 2 6-32 8-32 10-24 1 -20 4 5 -18 16 3 -16 8

STAINLESS STEEL (TYPE 304)

ALUMINUM ALLOY (5000 SERIES)

BRASS (70-30 AND 65-35)

MAXIMUM FASTENING TORQUE(A) n·m lbf · in. 0.7 6 1.4 12 1.6 14 4.9 43

ULTIMATE TENSILE LOAD kn lbf 2.2 500 3.4 765 4.3 960 7.8 1750

MAXIMUM SHEAR LOAD kn lbf 1.7 375 2.6 575 3.2 720 5.8 1300

8.1

72

12.9

2900

9.8

2200

11.9

106

19.1

4300

14.6

3250

1.1 2.3 2.6 8.5

10 20 23 75

3.5 5.6 6.8 12.8

790 1260 1530 2880

2.6 4.1 5.2 9.6

590 940 1150 2160

14.2

126

16.7

3750

12.5

2800

20.0

186

21.6

4850

16.0

3600

0.40 0.85 1.13 4.5

3.5 7.5 10 40

1.7 2.6 3.3 6.0

375 585 735 1360

1.0 1.6 2.0 3.8

235 365 460 850

7.9

70

10.2

2300

6.2

1400

9.1

81

15.1

3400

9.3

2100

15.8

140

25.0

5500

13.4

3000

0.9 1.8 2.09 6.9

8 16 18.5 61

2.7 3.8 4.6 8.7

600 860 1040 1950

1.7 2.5 3.0 5.7

390 560 680 1275

11.5

102

14.6

3280

9.5

2140

16.9

150

21.4

4800

14.1

3160

(A) THESE VALUES SHOULD DEVELOP FASTENER TENSION TO SLIGHTLY LESS THAN YIELD POINT AND SHOULD BE USED ONLY AS A GUIDE; BASIC SPECIFICATIONS SHOWN ABOVE COVER TYPICAL MECHANICAL PROPERTY VALUES.

The stud and workpiece materials can be common low-carbon steel, stainless steel, or aluminum. Also used are mediumcarbon steel, lead-free brass and copper, Inconel, titanium, René 41, zirconium, gold, silver, and platinum. Table 2 shows

some of the more popular combinations of materials. Various studs and plate materials can be combined, except when aluminum-base material is used.

TABLE 2 WELDING CAPABILITIES OF CAPACITOR DISCHARGE STUD FASTENERS BASE MATERIAL

MILD STEEL (1006 THROUGH 1030) MEDIUM-CARBON STEEL (1030 THROUGH 1050) GALVANIZED SHEET DUCT OR DECKING STRUCTURAL STEEL STAINLESS STEEL (TYPES 405, 410, 430, AND 300 SERIES, EXCEPT 303) LEAD-FREE BRASS, ELECTROLYTIC COPPER, LEADFREE ROLLED COPPER MOST ALUMINUM ALLOYS OF THE 1000, 3000, 5000, AND 6000 SERIES(B) DIE-CAST ZINC ALLOYS

STUD MATERIAL MILD STEEL STAINLESS (TYPES 1008, STEEL 1010) (TYPES 304, 305)

ALUMINUM (5356, 6061)

BRASS (70-30, 65-35)

EXCELLENT GOOD(A)

EXCELLENT GOOD(A)

... ...

EXCELLENT GOOD(A)

GOOD(A)

GOOD(A)

...

...

GOOD(A) EXCELLENT

GOOD(A) EXCELLENT

... ...

EXCELLENT EXCELLENT

EXCELLENT

EXCELLENT

...

EXCELLENT

...

...

EXCELLENT . . .

GOOD(A)

GOOD(A)

EXCELLENT GOOD(A)

(A) GENERALLY FULL-STRENGTH RESULTS, DEPENDING ON THE COMBINATION OF STUD SIZE AND BASE METAL. (B) OTHER MATERIALS, SUCH AS 7000 SERIES ALUMINUM, TITANIUM ALLOYS, INCONEL, AND SO ON, CAN BE WELDED UNDER SPECIFIED CONDITIONS. CD stud welding is used in many industries in a large variety of applications because it is one of the most versatile and reliable processes available. The food processing industry utilizes this process on stainless steel utensils and machinery, aluminum cookware, and appliances. The aerospace industry applies the process to aircraft engines, aircraft parts and instrument areas, and components for outer space vehicles. The appliance industry uses CD welding for its cabinets, trims, and other components. The building industry uses it for doors, facia, subassemblies, and other exposed finished areas. The insulation industry uses it in securing insulation to heating and air conditioning duct work, as well as on apparatus at power plants, schools, and other major buildings. The automotive industry uses the process to secure trim and various components to cars and trucks. The ship-building industry uses it in fabricating components in ship kitchens, in insulation applications aboard ship and in other areas where components and instruments must be secured. The electrical industry uses this process when securing components inside various controllers and equipment, as well as in trim parts for cabinets. The sign industry uses it on both structural and component members of large and small signs. Capacitor Discharge Stud Welding Richard L. Alley, American Welding Society

Equipment The equipment used for CD stud welding can be either portable or stationary production-type units. The portable units consist of the basic controller that utilizes standard 110 V alternating current (ac) power input to charge the capacitors and a lightweight gun-shaped tool used for placing and welding the fasteners. The production-type units typically require

240/480 V ac three-phase incoming power and a compressed air supply. The production units are normally used for higher rates of productivity, close-tolerance work (down to ±0.13 mm, or 0.005 in.), critical reverse side marking requirements, automatic stud feeding, automatic feeding of the part to be welded, automatic location of the stud on the part, and exotic materials. The production rate of portable equipment normally ranges from 4 to 6 studs/min, depending on the application. Material handling and other factors could cause this rate to be either slower or faster. The production rates for the permanent production equipment can range from 3 to 20 studs/min, depending on the application, special requirements, and tooling. When the initial-gap or initial-contact modes of stud welding are used, welds are made on all materials without the use of shielding gases or ceramic arc shields. When using the drawn-arc mode while working with aluminum, a shielding gas would be necessary. This mode is also less flexible when working with exotic materials. Capacitor Discharge Stud Welding Richard L. Alley, American Welding Society

Personnel Responsibilities The CD stud welding process is easy to operate and maintain. The operator must first ensure that his unit is connected to the proper incoming voltage. The operator then determines whether the weld gun is connected to the negative and the ground is connected to the positive. Straight polarity is used for welding in nearly all applications. The operator places a given diameter of fastener into the stud gun. The gun is then placed on the work surface. The operator closes the trigger of the gun and the welding process begins. After the welding process is completed, the operator lifts the gun off the welded stud. The operator needs to inspect the welds to ensure that the settings were correct. The controller voltage, which is set according to the recommendations of the equipment manufacturer, depends on the material and diameter of the fastener to be welded, as well as on the thickness and type of base material being used. There are also other parameters that should be considered. If it is desirable to reduce spatter and smoky conditions in the weld area, then wetting the weld area with a wetting agent just prior to making the weld is recommended. If a wetting agent is used, then it is best to add a few drops of liquid hand soap to a pint of plain water and to spray or brush the solution into the weld area. This will produce a clean weld with reduced, if not eliminated, spatter in the weld area. However, the noise of welding increases when a wet weld is made. Weld inspection is normally done by bending the fastener by applying either a torque or tensile force to it. If a tensile force is applied, then care should be given to the diameter of the hole used to secure the plate material. A hole that is as small as possible should be used to accommodate the flange of the stud. If a large hole is used, the tensile factors may be reduced, because the plate material could yield prematurely. Prior to production welding, it is recommended that sample welds be made on similar materials and tested to destruction. During production, operators should visually inspect the welds periodically. If weld inconsistency occurs, the process should cease, and the cause should be determined and eliminated. A common occurrence that contributes to inconsistency is movement of the plate material during the short weld cycle. This can occur when working with materials that are less than 3.2 mm (0.125 in.) thick. There should not be a gap between the workpiece and the backup material of the fixture or table being used. Proper fixtures and attention to their clamping system can alleviate this problem.

Capacitor Discharge Stud Welding Richard L. Alley, American Welding Society

Reference

1. R.L. O'BRIEN, ED., WELDING HANDBOOK, 8TH ED., VOL 2, AWS, 1991, P 317-323 Capacitor Discharge Stud Welding Richard L. Alley, American Welding Society

Selected References

• RECOMMENDED PRACTICES FOR STUD WELDING, C5.4-84, AWS, 1984 Plasma-MIG Welding Ian D. Harris, Edison Welding Institute

Introduction PLASMA-MIG WELDING can be defined as a combination of plasma arc welding (PAW) and gas-metal arc welding (GMAW) within a single torch, where a filler wire is fed through the plasma nozzle orifice. The process can be used for both welding and surfacing. The plasma-metal inert gas (MIG) welding process was invented at the Philips Research Laboratories in Eindhoven, Netherlands, around 1969 (Ref 1). For descriptions of the PAW and GMAW processes, see the articles "Plasma Arc Welding," and "Gas-Metal Arc Welding" in this Section of the Handbook. The principles of operation, in terms of equipment, are illustrated in Fig. 1. Separate power supplies are used for the PAW and the GMAW elements of the equipment. An arc is struck between the tungsten electrode and the workpiece in a similar fashion to that of a PAW system. The filler wire can be fed to the plasma arc, either with or without the GMAW arc established. Without power supplied to the filler wire, the system can be operated as a PAW system with concentric feed of filler wire. Later versions of the system incorporated an annular electrode to replace the offset tungsten electrode in the welding torch.

FIG. 1 SCHEMATIC OF PLASMA-MIG WELDING EQUIPMENT. SOURCE: REF 1

Current and Operating Modes. The equipment can be operated either with a single power source, effectively as a PAW system with concentric filler wire feed, or with two power sources, for the plasma-MIG operation.

The polarity of the tungsten electrode is direct current, electrode negative (DCEN), as is that of the GMAW part of the system. The heat of the plasma arc is sufficient to achieve good metal transfer stability for the GMAW element, despite the fact that when this process is used separately, it is almost always used in a direct-current, electrode positive (DCEP) mode. The filler wire is heated by the constricted plasma arc, as well as by the cathode heating of its own arc, and by resistance heating along the wire extension. Therefore, the melting and deposition rates of the wire are higher than the rates achieved by heating with either arc alone. Metal transfer is governed not only by plasma streaming, but also by arc forces between the wire tip and the workpiece. Because the metal droplets are totally enclosed by the plasma stream, spray transfer takes place even though the GMAW element operates on negative polarity. Advantages and Disadvantages. The advantages of the plasma-MIG process include deposition rates and joint

completion rates that are higher than those of the conventional GMAW process. The independent control of the plasma arc and current to the filler wire leads to more control of metal deposition. This capability can yield improved productivity and good flexibility for controlling heat input and arc characteristics in both welding and surfacing operations. Good control of dilution is achieved by running the system without any power applied to the filler wire. Metal transfer stability is increased, compared to that of the conventional GMAW process, and results in lower spatter levels. The cleaning action of the plasma arc results in lower porosity in aluminum alloys, compared to that of the conventional GMAW process. Disadvantages include the capital cost of two power sources (although there are systems that are designed to operate with one), the greater complexity of the torch, and the increased maintenance time and cost associated with this complexity. With two power sources, more welding parameters need to be set up, compared to the conventional GMAW process. Equipment As noted earlier, the basic equipment includes a power source for the plasma arc and a power source for the GMAW part of the system. A special torch incorporating both a contact tip for the GMAW element and a cathode for the PAW element is required. The initial design incorporated an offset tungsten electrode, as well as a concentric conduit and

contact tip for the delivery of the consumable wire (Fig. 1). A later design incorporated a concentric cathode for the plasma arc (Fig. 2).

FIG. 2 SCHEMATIC OF MODERN PLASMA-MIG TORCH WITH ANNULAR PAW ELECTRODE AND ADDITIONAL (FOCUSING) GAS STREAM. SOURCE: REF 2

The plasma-MIG torch can be readily fitted to existing welding equipmentsuch as side beams and welding carriages, to replace the GMAW process in mechanized welding operations. Power Sources. A constant-current power source with a high-frequency circuit to initiate the pilot arc is used for the plasma arc component of the system. The power source for the GMAW component can be used as a constant-voltage or a constant-current rectifier. Power sources have welding currents that typically range from 40 to 200 A for the plasma arc and from 60 to 300 A for the GMAW element at 100% duty cycle. However, equipment with welding currents up to 800 A is available and can be used for surfacing applications. Welding Torches. A special torch with a concentric cathode for the plasma arc and a concentric conduit and contact tip

for the delivery of the consumable wire is required (Fig. 2). A water-cooled copper alloy nozzle is used to constrict the arc and to form a collimated plasma jet that exits the nozzle orifice. A plasma orifice gas and a focusing gas from the same supply are used; the latter is delivered via channels between the plasma welding electrode and the constricting nozzle. The focusing gas results in greater arc constriction and arc stability and prolongs the life of the constricting nozzle by creating a boundary gas layer between the nozzle orifice and the plasma arc. The contact tip, plasma cathode, and the constricting nozzle are all directly water cooled to provide 100% duty cycle at the typical welding currents identified in the preceding section, "Power Sources." Shielding Gases. Three shielding gases are utilized: one for the plasma (orifice) gas, one to provide additional arc

constriction and arc stability, and one for supplementary shielding. The plasma gas and the focusing gas are usually argon, because an inert gas is required to prevent oxidation of the PAW electrode. The supplementary shielding gas can be argon, argon-oxygen, argon-carbon dioxide, or argon-hydrogen, depending on the nature of the workpiece being welded or, in the case of a surfacing operation, on the material being deposited. Argon is used when welding aluminum alloys, whereas argon-oxygen and argon-carbon dioxide are used when welding steels. Argon-hydrogen is used when welding stainless steels or when surfacing with them.

References

1. W.G. ESSERS, A.C.H.G. LIEFKENS, AND G.W. TICHELAAR, PLASMA-MIG WELDING, PROC. CONF. ADVANCES IN WELDING PROCESSES, THE WELDING INSTITUTE, 1971, P 216-219 2. J.D. SWART, PLASMA-MIG BOOSTS TANK TRAILER OUTPUT, WELD. DES. FABR., FEB 1983, P 54-55, 59 Plasma-MIG Welding Ian D. Harris, Edison Welding Institute

Procedure Process Operating Procedure. The plasma arc is ignited using a pilot arc in a fashion similar to that of a PAW system. The main arc is transferred from the electrode to the workpiece and the plasma jet passes through the nozzle orifice. The system can be operated in this way, with the concentric cold wire being fed through the axis of the torch. A higher melting rate is achieved when power is applied to the wire through the contact tip and when both arcs are run simultaneously. The higher energy imparted to the wire by the plasma arc results in an increased wire deposition rate (Table 1). In this operating mode, deposition rates higher than those typical of the GMAW process can be achieved.

TABLE 1 COMPARISON OF FILLER WIRE MELTING RATES FOR PLASMA-MIG WELDING WITH AND WITHOUT THE GMAW ARC Mild steel filler wire; plasma gas, argon; shielding gas, 89% Ar + 6% CO 2 + 5% O2 PLASMA ARC CURRENT, A

110 135 160 190 190 190

PLASMA ARC VOLTAGE, V

29 30 32 34 37 38

FILLER WIRE CURRENT, A

... ... ... ... 100 150

MIG-ARC VOLTAGE, V

DIAMETER OF FILLER WIRE

MELTING RATE OF FILLER WIRE

... ... ... ... 31 32

mm 0.9 0.9 0.9 0.9 1.2 1.2

g/min 22 28 33 40 85 130

in. 0.04 0.04 0.04 0.04 0.05 0.05

oz/min 0.78 0.99 1.16 1.41 3.0 4.6

Source: Ref 1 Inspection and Weld Quality Control. Inspection requirements are similar to those of other arc welding or surfacing operations. Visual, ultrasonic, and radiographic inspection techniques are most appropriate. The dual action of the GMAW and plasma arcs results in weld quality that is sometimes higher than that achieved by the GMAW process alone. This is particularly true for aluminum alloys, because the cleaning action of the plasma arc often results in reduced porosity.

Quality control requires monitoring the welding parameters for both power sources, as well as monitoring the wire feed. In addition, the condition of the nozzle orifice (that is, the wear and concentricity of the orifice) should be monitored. Troubleshooting. The relatively complex nature of the welding torch involves increased maintenance time. Erosion of

the copper alloy nozzle orifice will cause a change in the arc shape and will affect the weld profile. Therefore, the nozzle should be checked periodically.


1. W.G. ESSERS, A.C.H.G. LIEFKENS, AND G.W. TICHELAAR, PLASMA-MIG WELDING, PROC. CONF. ADVANCES IN WELDING PROCESSES, THE WELDING INSTITUTE, 1971, P 216-219 Plasma-MIG Welding Ian D. Harris, Edison Welding Institute

Applications Material Types. The plasma-MIG process is suitable for welding a wide variety of materials. The high heat energy

supplied by the plasma and gas-metal arcs makes the process suitable for high-melting-point materials, such as tungsten and molybdenum. The most common application is welding aluminum sheet and plate. Wear-resistant steels are used with the process in hardfacing applications. Austenitic stainless steels, such as types 308, 309, and 347, as well as nickel alloys, such as alloy 625, are used in cladding applications. Both solid and flux-cored wires (Ref 3) can be employed for welding and surfacing, although most applications involve solid wires. Industries. The plasma-MIG welding process has been used for the deposition of corrosion-resisting stainless steel and

nickel-base alloys in the offshore industry, for the general fabrication of silos and tank trailers made from aluminum alloys (Ref 2), and for hardfacing applications in the excavation equipment industry, as well as the dredging and offshore industries. Typical Components and Joints. The plasma-MIG process is suitable for welding of joints and for surfacing

operations. The wide range of heat inputs available by choosing how to apply current to the consumable electrode provides additional flexibility for surfacing operations, compared to the range of heat inputs available with an external wire feed using the GTAW/PAW or the conventional GMAW process. Single-V butt joints are commonly used for the plasma-MIG welding of plate. A full-penetration weld can be made in a single pass on a 9.5 mm (

3 in.) thick mild steel plate when operating the plasma and gas-metal arcs simultaneously. This 8

compares to three passes when just the PAW process is used. One would expect the same joint to require two or three passes for the conventional GMAW process. When butt welding aluminum alloy sheet and plate in thicknesses ranging from 4.6 to 6.1 mm (0.18 to 0.24 in.), singlepass welds can be made at travel speeds that are more than twice as fast as when the GMAW process is used alone (Table 2). Lower porosity levels are achieved with the plasma-MIG process because of the cleaning action of the plasma arc. When using mild steel and stainless steel, the plasma-MIG process can make butt welds comparable to those made with the GMAW process alone at almost twice the travel speed (Table 2).

TABLE 2 COMPARISON OF WELDING SPEEDS FOR GMAW AND PLASMA-MIG WELDING OF ALUMINUM AND STEEL PARAMETER

SHEET THICKNESS, MM (IN.) WIRE DIAMETER, MM (IN.) GMAW CURRENT, A PLASMA CURRENT, A WIRE FEED SPEED, M/MIN (IN./MIN) WELDING SPEED, M/MIN

MATERIAL ALMG4.5MN GMAW PLASMAMIG

PLASMAMIG

4.6 (0.18) 1.6 (0.063) 280-300 ... 8.1 (320)

4.6 (0.18) 1.6 (0.063)

5.1 (0.20) 1.6 (0.063)

200 200 10.9 (430)

0.4 (16)

1.1 (43)

STEEL GMAW

PLASMAMIG

280 200 10.9 (430)

4.1 (0.16) 1.2 (0.047) 250 ... 5.1 (200)

4.1 (0.16) 1.2 (0.047) 280 200 11.4 (450)

0.97 (38)

0.51 (20)

0.97(38)

(IN./MIN) Source: Ref 2

The hardfacing of dredging equipment with martensitic steel was made more productive by using the plasma-MIG process, rather than the conventional GMAW process. Higher deposition rates were achieved, resulting in reduced labor costs for the semiautomatic welding operation. Deposition rates of up to 20 kg/h (44 lb/h) can be achieved by plasmaMIG surfacing of mild steel with stainless steel flux-cored wires (Table 3).

TABLE 3 WELD SURFACING PARAMETERS FOR STAINLESS STEEL DEPOSITS ON MILD STEEL USING PLASMA-MIG WITH FLUX-CORED WIRES

UNDILUTE D WELD METAL PZ 6400 (347) SURFACIN G PZ 6410 (308) SURFACIN G PZ 6415 (309) SURFACIN G

MIG WELDIN G CURREN T, A

MIG WELDIN G VOLTAG E, V

PLASMA WELDIN G CURREN T, A

PLASMA WELDIN G VOLTAG E, V

WELDING SPEED

DEPOSITIO N RATE

INNER WIRE STICKOUT

mm/ mi

in./mi n

kg/ h

lb/h

m m

in.

300

39

150

44

110

4.4

10

22

35

200

31

150

40

80

3.2

6.6 14. 6

300

44

150

46

140

5.6

13

400

50

150

51

190

7.6

20

DILUTIO N, %

FERRIT E NO.

1. 4

...

15, 12, 13

55

2. 2

8.6

12

28. 7

55

2. 2

13.8

5.5

44

55

2. 2

9.7

6.5

Source: Ref 3


2. J.D. SWART, PLASMA-MIG BOOSTS TANK TRAILER OUTPUT, WELD. DES. FABR., FEB 1983, P 54-55, 59 3. F. EICHHORN AND E. VAN GAEVER, ADVANTAGEOUS SURFACING WITH PLASMA-MIG USING CORED WIRES, PROC. 1ST INT. CONF. SURFACE ENGINEERING, THE WELDING INSTITUTE, 1986, P 99-110 Plasma-MIG Welding Ian D. Harris, Edison Welding Institute

Personnel Skill Level and Training. The plasma-MIG process can be used for semiautomatic, mechanized, or automated operation. It requires a skill level comparable to that of the conventional GMAW process, although additional parameters

need to be set for the plasma arc. Most welding and surfacing activities are carried out with mechanized equipment. When welding parameters are set, operation is similar to that of the GMAW process. Health and safety issues are similar to those of other arc welding processes. Electric shock, eye protection, burns, ultraviolet radiation, and fume exposure are typical concerns. The plasma-MIG process generates more radiant heat than the conventional GMAW process. Plasma-MIG Welding Ian D. Harris, Edison Welding Institute

References

1. W.G. ESSERS, A.C.H.G. LIEFKENS, AND G.W. TICHELAAR, PLASMA-MIG WELDING, PROC. CONF. ADVANCES IN WELDING PROCESSES, THE WELDING INSTITUTE, 1971, P 216-219 2. J.D. SWART, PLASMA-MIG BOOSTS TANK TRAILER OUTPUT, WELD. DES. FABR., FEB 1983, P 54-55, 59 3. F. EICHHORN AND E. VAN GAEVER, ADVANTAGEOUS SURFACING WITH PLASMA-MIG USING CORED WIRES, PROC. 1ST INT. CONF. SURFACE ENGINEERING, THE WELDING INSTITUTE, 1986, P 99-110 Resistance Spot Welding Neville T. Williams, British Steel

Introduction RESISTANCE SPOT WELDING (RSW) is a process in which faying surfaces are joined in one or more spots by the heat generated by resistance to the flow of electric current through workpieces that are held together under force by electrodes. The contacting surfaces in the region of current concentration are heated by a short-time pulse of low-voltage, highamperage current to form a fused nugget of weld metal. When the flow of current ceases, the electrode force is maintained while the weld metal rapidly cools and solidifies. The electrodes are retracted after each weld, which usually is completed in a fraction of a second. Spot welding is the most widely used joining technique for the assembly of sheet metal products such as automotive body-in-white assemblies, domestic appliances, furniture, building products, enclosures and, to a limited extent, aircraft components. Many assemblies of two or more sheet metal stampings that do not require gas-tight or liquid-tight joints can be more economically joined by high-speed RSW than by mechanical methods. Containers frequently are spot welded. The attachment of braces, brackets, pads, or clips to formed sheet-metal parts such as cases, covers, bases, or trays is another common application of RSW. Major advantages of spot welding include high operating speeds and suitability for automation or robotization and inclusion in high-production assembly lines together with other fabricating operations. With automatic control of current, timing, and electrode force, sound spot welds can be produced consistently at high production rates and low unit labor costs using semiskilled operators. Most metals can be resistance spot welded if the appropriate equipment is used coupled with suitable welding conditions. This is particularly true for thin sheet or strip steel products, whether uncoated or coated. Resistance Spot Welding

Neville T. Williams, British Steel

Equipment The equipment needed for RSW can be simple and inexpensive or complex and costly, depending on the degree of automation. Spot welding machines are composed of three principal elements: • • •

ELECTRICAL CIRCUIT, WHICH CONSISTS OF A WELDING TRANSFORMER, TAP SWITCH, AND A SECONDARY CIRCUIT CONTROL CIRCUIT, WHICH INITIATES AND TIMES THE DURATION OF CURRENT FLOW AND REGULATES THE WELDING CURRENT MECHANICAL SYSTEM, WHICH CONSISTS OF THE FRAME, FIXTURES, AND OTHER DEVICES THAT HOLD AND CLAMP THE WORKPIECE AND APPLY THE WELDING FORCE

Specifications for resistance welding equipment have been standardized by the Resistance Welder Manufacturers Association (RWMA), and specifications for controls are issued by the National Electrical Manufacturers Association (NEMA). Resistance Spot Welding Neville T. Williams, British Steel

Electrical Circuit The transformer used in a direct energy resistance delivers the power to the workpiece by changing the input highvoltage, low-amperage, alternating current (ac) in the primary winding to a low-voltage, high-amperage current in the secondary winding. This system forms the basis of pedestal, or gun-type, welding machines where the output of the secondary transformer is applied directly to the welding electrodes. More complex secondary circuits are used in multiwelders. To minimize the size and cost of multiwelders, transformers can be designed with two or more secondary circuits supplied from the same primary circuit. This design is achieved by using transformers in which the primary current is supplied to two or more separate secondary windings. Split or multiple secondary transformers are generally used to produce four or more welds at the same time. One controller is used for control of the timing and current supplied to the electrode, which means that it is not possible, under these circumstances, to achieve individual control or setting of the welding conditions for each pair of welding electrodes. Difficulties can arise in optimizing the necessary welding conditions to maximize the electrode life for each welding station under these conditions. In addition, the possibility exists for developing shunt currents, or an out-of-balance welding current, between each secondary circuit. Multiwelders make large use of single-welding and multiple-welding configurations based on direct, indirect, series, and push-pull secondary circuits (see Fig. 7 of the article "Procedure Development and Practice Considerations for Resistance Welding" in this Volume). Single-Weld Configurations. Most pedestal and portable gun-welding machines employ direct welding, where the

current flows directly through the welding electrodes and the workpiece to produce a single weld. In areas of difficult access, indirect welding may be necessary. Indirect welding is generally used when welding is carried out using either stationary or indexing welding guns and when access to the sheets being welded is limited to one side. In this case, the backing electrode generally consists of a large bar or platen, and the welding current flows from one side of the transformer secondary to the active electrode, through the welded area, and then along the platen/bar electrode or connecting jumper cable to the dummy electrode, which is connected to the other leg of the transformer secondary. Higher welding currents are generally necessary when using dummy welding guns, which can lead to a shorter electrode life. The extent to which larger currents are required depends largely on the distance between the welding electrode and the dummy return electrode. In a multiwelder, this distance can be either fixed or variable, depending on whether a

stationary or indexing head is used. Both types of welding heads can lead to losses due to stray current paths and extensive shunting of the current through jigs and fixtures. These factors need to be considered in any design concept for a multiwelder station. Two-Weld Configurations. Because of the need to minimize the number of transformers and the overall size and cost of

a multiwelder, it is necessary to utilize configurations whereby two welds are made in one stroke of an electrode. Parallel, series, or push-pull welding configurations are used for this purpose. In parallel welding, two or more direct welds, usually closely spaced, are made simultaneously with all upper

electrodes connected to one side of the transformer, and all bottom electrodes connected to the other side of the transformer. Parallel welding has the advantage of lower shunt currents, thereby allowing shorter interweld spacings to be used. However, if two or more circuits are fed from a common transformer secondary winding, the current in each circuit will not be equal unless the impedances of each circuit are identical. This is often difficult to achieve in practice, and the electrode life obtained can, therefore, be different for each of the electrode pairs connected to the same transformer. Series and Push-Pull Welding. Two simultaneous welds can also be produced by means of series welding in which the

secondary circuit is connected to two contoured electrodes that contact the workpiece from the same side, and a current backup bar/platen is used on the reverse side. Push-pull welding is also employed whereby either one or two welds are made, depending on the electrode configuration. In push-pull RSW, two transformers are used to form the circuit between the welding electrodes; such a configuration has advantages, particularly where a deep-throated machine is required. Such machines ordinarily require a large secondary loop because of the need for conducting the welding current around the large working area. Using a push-pull configuration, where the upper set of transformers matches the lower set, eliminates the need for long cable loops to the electrodes, resulting in a lower secondary impedance. A longer electrode life can be obtained in push-pull welding compared to series welding. It should also be noted that current shunting between the electrodes occurs in both series and push-pull welding configurations, but to a lower extent in the latter. The magnitude of the shunt currents depends on interelectrode distance and the type of material being welded. In all cases, current shunting will lead to a lower electrode life when welding zinccoated steels due to the need to use higher welding currents to compensate for the shunt currents (this can also lead to additional electrode sticking problems) and to the rapid falloff in weld size that occurs under conditions where current compensation is insufficient or nonexistent. To some extent, the effects of both high secondary circuit impedance and current shunting can be overcome by using a constant-current facility fitted into the welding machine control circuit, but there is a limit to the effectiveness of these controls in a production situation. In addition, direct current (dc) welding systems can be used in which secondary inductance/reactance effects are minimized. However, the effectiveness of using dc depends on the method used to develop the dc waveform and on the extent of the superimposed ac ripple on the dc waveform. Secondary Impedance. The electrical characteristics of a spot welding machine are best defined in terms of their effect on the impedance of the secondary circuit. The latter is composed of two components: •

•

REACTANCE, WHICH IS DETERMINED PRIMARILY BY THE DEPTH AND HEIGHT OF THE THROAT OF THE MACHINE AND THE AMOUNT OF STEEL IN THE THROAT. ANOTHER CRITICAL FACTOR IS THE FREQUENCY OF THE APPLIED CURRENT. RESISTANCE, WHICH DEPENDS ON THE LENGTH AND CROSS SECTION OF THE CONDUCTORS. CONDUCTOR MATERIAL AND OPERATING TEMPERATURE NEED ALSO TO BE CONSIDERED.

Cables can be classified into three types: air-cooled jumper cables, water-cooled jumper cables, and kickless cables. In any application, it is necessary to ensure that the type and dimensions of cables are chosen to achieve the best balance between the total circuit resistance and the electrode life to be obtained with a multiple spot welder. Resistance Spot Welding Neville T. Williams, British Steel

Machine Construction Welding operations in highly automated production lines are based primarily on multiple spot welders and robotic cells. In addition, manual welding operations can be used to manufacture either subassemblies, which are fed into the main production/assembly lines, or, in many instances, finished products. These differing end uses require welding machines of varying designs and characteristics. RSW machines can be divided into three basic types: • • •

PEDESTAL-TYPE WELDING MACHINES PORTABLE WELDING GUNS MULTIPLE SPOT WELDING MACHINES INCORPORATING LIGHTWEIGHT GUN WELDING UNITS

Pedestal-type welding machines form the basis of most resistance spot or projection welding operations. Standard

machines are either of the rocker arm type, as determined by the rocker action of the moving upper electrode arm, or stationary, direct-acting machines. The latter outnumber the former because rocker arm machines are difficult to align due to the movement of the electrode in an arc motion. As the electrode wears, various skidding actions can be created between the upper and lower welding electrodes, depending on the extent of the electrode wear. This can change the point at which the electrodes meet relative to the fulcrum point of the moving arm. If the electrodes meet at a point above the fulcrum point, they have a tendency to slide outward from the machine. On the other hand, if the electrodes meet on a line below the fulcrum of the moving arm, the electrodes show a tendency to slide inward toward the machine. These relative movements can be different between the top and bottom electrodes, a consequence of which is a reduction in electrode life, particularly if welding coated steels. Other advantages of direct-acting machines include a lower electrode assembly mass, faster "follow up" during welding, and a degree of rigidity that is impossible to obtain with an ordinary rocker arm machine. Press-type direct-acting machines have an upper electrode and welding head that move vertically in a straight line. The welding head is guided in bearings or guideways of sufficient proportion to withstand the offset loads imposed upon them. All of these factors tend to increase the electrode life obtained in a particular welding cell. Portable welding guns are used in RSW when it is impractical or inconvenient to transport the work to the machine.

The portable welding gun consists of water-cooled electrode holders, an air or hydraulic actuating cylinder, hand grips, and an initiating switch. The gun usually is suspended from an adjustable balancing unit, and the welding force is supplied by an air cylinder. Because of the high secondary losses of portable machines, transformers used in these machines have a secondary voltage two to four times as great as the voltage of transformers used in stationary machines of equal rating. Welding current is transmitted between the transformer terminals and gun terminals through a secondary cable, usually of the low-impedance or kickless type. The reactance of this type of cable is near zero, which results in a high power factor and reduced kilovolt-ampere (kVA) demand. For safety reasons, the current initiating switch is usually operated at low voltage. Recent developments in encapsulated transformers and high-speed fail-safe devices enable the transformer to be integral with the welding gun. This results in lower secondary losses and is particularly preferred in robotic applications. Gun welders typical of portable manually operated and robotic welding cells can be classified into two basic types: • •

S-TYPE (SCISSORS) J-TYPE (DIRECT ACTION OR C/J TYPE)

Two important features that influence the performance of these types of welding guns are the rate of force buildup and the extent to which the arms of the gun deflect under load. These features characterize the welding gun in terms of rigidity and the mechanics involved in their operation. An S-type gun utilizes a lever action to transmit the force applied by the cylinder head to the welding electrodes. However, due to an unbalanced pivot point, the force measured at the electrodes is less than that applied at the cylinder. With J-type welding guns, the force applied by the pneumatic cylinder is transmitted directly to the welding electrode. The rigidity of the framework of the welding gun is characterized by the extent to which the electrode arms deflect on application of the electrode force. This is an important factor in determining weld quality and the amount of skidding that can occur between the two welding electrodes. It should be noted that extensive skidding can lead to elliptical welds and/or decreased electrode life.

Resistance Spot Welding Neville T. Williams, British Steel

Welding Electrodes Materials for RSW electrodes should have sufficiently high thermal and electrical conductivities and sufficiently low contact resistance to prevent burning of the workpiece surface or alloying at the electrode face. In addition, the electrode should have adequate strength to resist deformation at operating pressures and temperatures. Because the part of the electrode that contacts the workpiece becomes heated to high temperatures during welding, hardness and annealing temperatures must also be considered. Electrode materials for RSW have been classified by RWMA and in International Standards Organization (ISO) standard ISO 5182. When welding two sheets of thickness up to 3 mm (

1 in.) using truncated cone-type electrodes, the electrode tip diameter 8

should be determined from:

D2 = t

(EQ 1)

where d2 is the initial tip diameter (in mm) and t is the thickness (in mm) of the sheet in contact with the electrode. When using truncated cone electrodes, the initial or set-up weld diameter should be equal to the diameter of the electrode tip:

D = D2 = 5 t

(EQ 2)

where d is the weld diameter (in mm). Whenever welding two sheets of dissimilar thicknesses, the electrode dimensions and the required weld size should be specified appropriate to the thinner sheet thickness. In the case of three thicknesses, the second thinnest sheet should be used. When using domed or pointed electrodes, Eq 1 may not apply and the electrode dimensions will depend on accessibility and flange width. In this case, the electrode tip dimensions and welding conditions should be chosen to give an initial weld diameter as specified in Eq 2. Where a pad or mandrel is used as the second electrode, its surface must be maintained to match the profile of the work. During normal production, electrodes tend to mushroom leading to an increase in electrode tip size. According to trends in automotive practice and international standards, the diameter of at least one of the electrode tips should not normally be allowed to increase above a value that results in a reduction greater than 30% from the starting weld diameter during a production run. This is equivalent to an increase in the electrode diameter to 1.3d. When this diameter has been reached, the electrode should be replaced or redressed to its original size and contour. Electrode shapes have been standardized by RWMA. Figure 1 shows the six standard face or nose shapes, identified by

letters A through F. Electrodes with type A (pointed) tips are used in applications for which full diameter tips are too wide. Type D (eccentric, formerly called offset) faces are used in comers or close to upturned flanges. Special tools are available for dressing the electrode faces, either in or out of the welding machine.

FIG. 1 STANDARD-RESISTANCE SPOT WELDING ELECTRODE FACE AND NOSE CONFIGURATIONS

Electrode Coolant Parameters. Generally, it is recommended that the water flow to the electrodes should be a 1 minimum of 4 L/min (1.1 gal/min) for welding two uncoated steel sheets of thickness up to and including 3.0 mm ( in.). 8

Higher flow rates are recommended when welding coated steels. The internal water cooling feed tube should be arranged to ensure that the water impinges onto the back working face of the electrode. The distance between the back and the working face of the electrode should not exceed the values given in the appropriate ISO or RWMA standard. The inlet water temperature should not exceed 20 °C (70 °F), and the outlet temperature should not exceed 30 °C (85 °F). To maintain these temperature levels, the electrode cooling water supply should be independent of transformer and thyristor water cooling circuits. Separate water circuits should be used for both top and bottom electrodes. Resistance Spot Welding Neville T. Williams, British Steel

Weldability of Steel Uncoated Steels. Low-carbon steel can be satisfactorily resistance spot welded using a wide range of time, current, and electrode force parameters. Practices recommended for spot welding SAW 1010 steel have been issued by the RWMA (Ref 1) and the American Welding Society (Ref 2). These data can serve as starting points for the establishment of optimum settings but may require modification, depending on the type of welding machine, the dynamic properties of the welding machine, the pneumatic characteristics, the characteristics of the secondary circuit, the electrode shape, and the material. Sufficient squeeze time should be chosen to enable the electrode force to build up to its preset value. For lightweight gun welders with limited force capability, the electrode force values are reduced up to 30% for sheet

thicknesses greater than 1.6 mm (

1 in.). The welding current needs to be adjusted accordingly. When welding sheets of 16

dissimilar thickness, welding conditions may be based on the thinner sheet, or on the second thinnest sheet when welding three thicknesses. In the case of high-strength low-alloy (HSLA) steels, up to 20% higher electrode force may be necessary. Welding currents may be reduced up to 20%, depending on the type of high-strength steel being welded. Pulse welding schedules may be necessary at the higher carbon equivalent levels and at sheet thicknesses greater than 2.0 mm (

5 in.). The 64

welding conditions required may depend greatly on the analysis of the steel because of differences in the resistivity of the various types of high-strength steel. For example, carbon content has the greatest effect on weldability of steels; weld

hardness increases rapidly with small increases in carbon content. This high hardness causes nugget interfacial tears and nugget deterioration. To obtain acceptable weld performance, carbon content should be kept below 0.10% + 0.3t, where t is the sheet thickness in inches. For materials above this range, postweld tempering may be necessary. Low-carbon steels are generally considered to be spot weldable. These steels can be obtained from a number of suppliers with reasonable certainty of only minor variations in weldability. On the other hand, HSLA steels are sold on the basis of strength, with each steelmaker producing its own composition. Therefore, this variation in composition may cause variations in weldability, which the steel user must consider. Specific alloying elements or combinations of elements can impart desired nugget properties in low-carbon or HSLA steels. The effects of particular elements are discussed below. Phosphorus and Sulfur. Phosphorus and, to some extent, sulfur, are generally considered to promote nugget interface

tearing. When the total content of phosphorus, sulfur, and carbon exceeds a critical value, spot weld interfacial failure is observed. This interfacial tearing during peel testing results in reduced current ranges and minimal lobe curve areas. Titanium. Data on the effect of titanium on both hot-rolled and cold-rolled steel indicate that the spot weldability of these

steels is not as good as that of the HSLA steels containing niobium and/or vanadium. Increasing titanium content generally reduces maximum button diameter, shear and cross-weld tensile strength, and welding current range. Titanium content should not exceed 0.18%, and the use of oversized electrodes and higher force may be required when welding titanium-bearing steels. Nitrogen. In HSLA steels, nitrogen promotes interfacial nugget failure. Nitrogen appears to be more critical in unkilled

cold-rolled steel. However, the sensitivity can be reduced by decreasing the nitrogen content or by combining the nitrogen with aluminum as in aluminum-killed steels. Zinc-Coated Steels. The present trend in the automotive industry toward the use of larger amounts of zinc-coated steels

in automotive assemblies demands that certain strict guidelines regarding the selection of equipment, choice of welding schedules, and maintenance procedures are rigidly followed. Two basic types of alloy coating are available: • •

ZINC-IRON (BOTH HOT-DIP AND ELECTRODEPOSITED) ELECTRODEPOSITED ZINC-NICKEL (SUPPLIED WITH OR WITHOUT AN ORGANIC PRIMER TOP COAT)

The available welding ranges for zinc-iron and zinc-nickel coated steels are similar to those for uncoated mild steel although displaced toward slightly higher currents (Fig. 2). When these steels are coated, it is essential that the film thickness not exceed 1 to 1.5 μm (40 to 60 μin.) in order to facilitate breakthrough of the film to enable current flow between the welding electrodes at the low secondary voltages typical of RSW machines.

FIG. 2 PLOT OF WELD TIME VERSUS SECONDARY WELD CURRENT TO OBTAIN WELDABILITY LOBES FOR SELECTED 0.8 MM (0.03 IN.) THICK STEELS. ELECTRODE PARAMETERS: FORCE, 1.8 KN (0.20 TONF); TIP DIAMETER, 5.0 MM (0.20 IN.)

Welding conditions can be optimized easily when using pedestal-type welding machines; however, this may not be the case in body-in-white assembly where a combination of multiple spot welders and robotic cells are more commonly used. Such operations are susceptible to poor fit-up of parts, access problems, and electrode configurations that give rise to electrode skidding. Consequently, the welding parameters may need to be adjusted to compensate for such situations. In many instances, one welding gun is called on to weld a range of sheet thickness combinations or even a variety of material combinations in a particular part. Although a multi-program welding control can be used to adjust for the different welding currents necessary for the various combinations, the actual current values may not be optimum because, more often than not, a single compromise electrode force value has to be used. In these circumstances, the electrode life obtained may be lower than that derived using optimized welding conditions. For example, the electrode life obtained can depend greatly on the type of welding gun used (Fig. 3) or the angle of approach of the electrode. For multiple spot welders, it is essential that all welding stations are balanced electrically and that the same air pressure is supplied to each station. Electrode shape and configuration, water cooling arrangements, and electrode dressing schedules need to be optimized. The last item is most important because satisfactory weld quality can only be provided when the necessary electrode condition and shape are maintained over a production run.

FIG. 3 EFFECT OF RSW ELECTRODE GUN TYPE ON ELECTRODE LIFE FOR SELECTED STEELS AND ALLOYS


1. RESISTANCE WELDING MANUAL, 4TH ED., RESISTANCE WELDER MANUFACTURERS ASSOCIATION, 1989 2. RECOMMENDED PRACTICES FOR AUTOMOTIVE WELD QUALITY--RESISTANCE SPOT WELDING, ANSI/AWS D8.7-88, AMERICAN WELDING SOCIETY, 1988 Resistance Spot Welding Neville T. Williams, British Steel

Weldability of Aluminum Alloys Most of the commercial aluminum alloys produced in sheet form can be spot welded provided that suitable welding equipment is used. Because of the low electrical resistance of aluminum alloys, considerably higher welding currents are required relative to welding mild steel. In addition, the narrow plastic range between softening and melting means that welding forces, time, and current need to be closely controlled. For consistent weld quality, it is essential that the tenacious surface oxide films are removed by mechanical or chemical techniques prior to welding. To minimize the likelihood of cracking or porosity in the weld, careful control of the electrode force is necessary. If the force is applied too soon, the welding current may be insufficient due to a premature loss of contact resistance. Thus, correct initiation of the force relative to the current is essential. Dual force schedules are frequently used; therefore, a lowinertia welding head can be beneficial because it allows a more effective buildup of force to its final level. Because of the higher secondary currents required when welding aluminum alloys, direct current secondary rectified or frequency converter welding machines are favored, particularly if high-quality welds are required (such as in aircraft applications).

Resistance Spot Welding Neville T. Williams, British Steel

References

1. RESISTANCE WELDING MANUAL, 4TH ED., RESISTANCE WELDER MANUFACTURERS ASSOCIATION, 1989 2. RECOMMENDED PRACTICES FOR AUTOMOTIVE WELD QUALITY--RESISTANCE SPOT WELDING, ANSI/AWS D8.7-88, AMERICAN WELDING SOCIETY, 1988 Projection Welding Jerry E. Gould, Edison Welding Institute

Introduction PROJECTION WELDING (PW) is a variation of resistance welding in which current flow is concentrated at the point of contact with a local geometric extension of one (or both) of the parts being welded. These extensions, or projections, are used to concentrate heat generation at the point of contact. The process typically uses lower currents, lower forces, and shorter welding times than does a similar application without the projections. Projection welding is often used in the most difficult resistance-welding applications because a number of welds can be made at one time, which speeds up the manufacturing process. PW applications are generally categorized as being either embossed-projection welding or solid-projection welding. These variations are shown in Fig. 1 and 2. Embossed-projection welding is generally a sheet-to-sheet joining process, in which a projection is stamped onto one of the sheets to be joined. Then, resistance welding is conducted on a stack of sheets. Heat initially concentrates at the contact point and in the walls of the projection during resistance welding. Early in the process, the projection itself collapses back into the original sheet. However, the initial heating raises the local resistivity of the joint area, allowing continued resistance heating at this location. Weld development then proceeds in the conventional manner, by forming a fused weld nugget.

FIG. 1 TYPICAL STACKED CONFIGURATION FOR EMBOSSED-PROJECTION WELDING OF SHEET

FIG. 2 TYPICAL CONFIGURATION FOR SOLID-PROJECTION WELDING

Solid-projection welding requires that the projection be forged onto one of the two components. Then, during resistance welding, the contact point and the projection itself experience preferential heating. In this case, the projection cannot simply collapse, as it does in embossed-projection welding. Rather, the projection collapses by penetrating the opposing material and by extrusion to the periphery. When compared with embossed-projection welding, the resulting joints are solid-state, rather than fusion, welds. The actual joints are caused by a combination of material forging and diffusion bonding, much like they are in resistance butt and flash butt welding. Projection Welding Jerry E. Gould, Edison Welding Institute

Applications Examples of projection welding are shown in Fig. 3. These applications, which range from sheet-to-sheet joints, to crosswire welds, to annular attachments, to nut welds, to weld screws, include both embossed- and solid-projection types of welding. Specific welding examples, as well as material effects, are described below.

FIG. 3 DIFFERENT PROJECTION-WELDING CONFIGURATIONS

Embossed Projections Typical embossed-projection weld configurations are shown in Fig. 3(a), (b), (d), and (h). The embossed-projection technique is primarily used for sheet-to-sheet applications. It does involve the additional expense of stamping the projections and ensuring that the electrodes are properly located over the projections. However, the ability of the projection to create highly localized joints allows this technique to extend the applications of resistance welding. Most notably, projection welding widely increases the allowable thickness mismatch associated with resistance welding. Generally, resistance welding is considered adaptable to thickness variations of less than 3:1. However, with the addition of the projection, there is no limit to the difference in joinable thicknesses. Similarly, embossed-projection welding is highly adaptable to the joining of multiple-sheet stacks. The approach is similar in that the projections are used to localize heat at critical interfaces in the overall joint. Another more recent use of embossed-projection welding is in applications where external marking must be minimized. This is a concern when a primer has been applied to one side of one of the sheets being joined. Because it is very effective at localizing heat at the interface, and because it uses shorter welding times, the PW process is more adiabatic and there is less residual heat that can damage outside surfaces. However, surface marking still does occur.

Solid Projections A considerably wider range of solid-projection welding processes is commonly used in production applications. Annular projection welding, like embossed-projection welding, is commonly used to provide a highly localized joint and to minimize thermal damage to other parts of the structure. Common variations of solid-projection welding are described below. Annular Projection Welds. One of the most common applications of solid-projection welding is to attach either tubular

components or members with circular bases to flat or curved substrates. This is commonly accomplished by annular projection welding, in which the projection is machined onto the circular base or tubular section. Then, resistance welding is conducted to set the projection into the substrate. A typical weld geometry is shown in Fig. 3(g). Annular projection welding generally provides high-integrity joints and can be used for leak-tight applications. Cross-wire welding is a variation of solid-projection welding in which the projection is formed by the contact point of

two crossing wires. The geometry for cross-wire welding is shown in Fig. 3(n). Upon resistance welding, heat is maximized at the location of the wire-wire contact, and the parts are subsequently forged together. Depending on the application, either highly localized joints (minimal forging) or heavily forged joints can be made. In some applications, cross-wire welding is conducted at rates up to 30 welds/s. Weld nuts represent an application of solid-projection welding in which the projections themselves occur as actual extensions of material from the nut. Typically, weld nuts contain 3, 4, or 6 of these projections, or feet, spaced around the nut periphery. During welding, all projections are attached simultaneously. Examples of weld nut geometries are shown in Fig. 3(k) and 3(m). Edge-to-sheet welds are typically a cross between embossed- and solid-projection welding. They are used to attach the

end of a sheet to the flat face of an opposing sheet (Fig. 3f). The projections for welding are generally stamped into the face of the attaching sheet. During welding, the projection does collapse locally back into the base sheet. However, bonding is strictly solid state, similar to that of other solid-projection welding processes. Material Effects Solid-projection welds are essentially strain-assisted diffusion bonds. Because the projections typically collapse at very high temperatures (generally, within several hundred °C of the melting point), diffusion bonding can occur within the very limited available time (usually, less than 1 s). Not surprisingly, some of the material-related factors that affect diffusion bonding also affect solid-projection welding. The most notable of these is the ability of the metal to dissolve its own oxide. As a result, materials that do not energetically favor the solid solubility of oxygen, rather than the formation of the oxide at aluminum forging temperatures, will be relatively difficult to projection weld. The strength-temperature relationship also affects projection weldability. Materials that maintain their strengths at relatively high temperatures permit substantial heating to occur before projection collapse. This heat then becomes available to promote diffusion after projection collapse. Premature collapse results in lower temperatures in which diffusion can occur and in reduced current density, which prohibits further resistance heating. Bulk resistivity also plays a role in projection welding, but to a lesser degree. Increased bulk resistivity can reduce the effectiveness of the projection as a current concentrator. With increasing bulk resistivity, the tendency is for delocalized heating and general, rather than local, collapse of the projection. As a result, high-resistivity materials are more difficult to projection weld. Mild steels and low-alloy, nickel-base alloys are ideal materials for projection welding, because they readily dissolve their own oxides and have adequate strength-temperature and resistivity properties. Stainless steels and higher-alloy nickelbase materials become slightly more difficult to weld, because of the formation of more-stable chromium and aluminum oxides, increased high-temperature properties, and higher resistivities. Projection welding is commonly applied to copper and copper-base alloys. In many applications, projections are virtually required for resistance welding, because of the high conductivity of these materials.

On the other hand, aluminum and aluminum-base alloys are very difficult to projection weld. The aluminum oxide is so tenacious that solid-projection welding, in particular, is nearly impossible. In addition, it is very difficult to localize heat, because aluminum alloys soften at such low temperatures. Conventional titanium alloys are also relatively difficult to projection weld. Although titanium readily dissolves its own oxide, its high resistivity and low forging temperature generally cause premature collapse of the projection. Projection Welding Jerry E. Gould, Edison Welding Institute

Equipment Resistance projection welding generally utilizes slightly modified resistance-welding equipment. Because the nature of projection welding is to direct current flow to the desired weld locations, there is considerably more flexibility in tooling. A specific example is spot welding that requires electrodes of specific sizes and shapes in order to localize current for welding. Projection welding is typically done with large, flat electrodes. In most applications, tooling is simply shaped to match the contour of the part in the contacting location. Given the current-locating characteristic of projections, multiple projection welds can be made simultaneously. Projection welding does require high compliance or the rapid response of the loading system. This is necessary because projection welding requires the collapse of a projection, which, in turn, requires some motion of the welding head. If the follow-up of the welding head is insufficient, then a local loss of force will occur, potentially causing catastrophic overheating of the contacting surfaces. Therefore, PW systems typically utilize "fast follow-up" heads (Fig. 4), which consist of relatively low inertia (mass) head components and mechanisms to maintain a nominally constant force. The head shown in Fig. 4 uses a relatively low volume diaphragm assembly and a spring to maintain this force on the components. Most systems, however, use either these low-volume diaphragm assemblies or springs to accomplish fast follow-up.

FIG. 4 TYPICAL FAST FOLLOW-UP (LOW INERTIA) HEAD FOR PROJECTION WELDING. (A) IN OPEN POSITION. (B) IN POSITION FOR SQUEEZING AND HEATING THE PROJECTION. (C) AT INSTANT OF PROJECTION COLLAPSE AND START OF NUGGET FORMATION

Projection Welding Jerry E. Gould, Edison Welding Institute

Personnel Like all resistance welding processes, PW processes utilize basically automated equipment, which means that the expertise level to operate the equipment is relatively low. Generally, the skill levels and the health and safety issues are identical to those of conventional resistance spot welding. An area in which expertise is beneficial to hardware function and on-stream weld quality is operator attention to detail. Resistance welding is subject to variations in setup conditions, and projection welding may be extraordinarily sensitive to setup variations in situations where the placement of projections is critical. With limited expertise, operators or supervisory staff can visually inspect the characteristics of the process and the welds themselves to identify improper setups, as well as maintenance needs. These quality assurance efforts, which fall under the auspices of total quality management or total quality joining, can be established with a minimum of on-line training and experience. Projection Welding Jerry E. Gould, Edison Welding Institute

Process Requirements The types of hardware required for projection welding are analogous to those required for other resistance-welding processes. The embossed-projection welding of heavy-gage, intermediate-gage, and thin-gage sheet, as well as joints of

dissimilar thicknesses, is described below. Heavy-Gage Sheet Steels. The projection welding of heavy-gage steels, which involves an embossed-projection

welding process, has many of the characteristics of solid-projection welding particularly in its early stages, because of the thickness of the sheet. In addition to defining the geometry of the projection, projection designs should provide an annular relief for projection material that is forged to the side during welding. Projection and die geometries for steels that range from 3.12 to 6.22 mm (0.123 to 0.245 in.) are shown in Table 1. The process requirements for forming these welds are given in Table 2. All welding schedules described in this article are single-pulse welding schedules.

TABLE 1 PROJECTION AND DIE GEOMETRIES FOR WELDING A RANGE OF HEAVY-GAGE STEELS Data are for tool steel hardened to 50-52 HRC.

USS GAG E NO.

11 10 9 8 7 6 5 4 3

MATERIAL THICKNES S (T )

PROJECTION HEIGHT (H DIAMETER ), (D ), ±5% ±2%

mm

in.

mm

in.

mm

in.

mm

in.

mm

in.

mm

in.

DIE HOLE DIAMETER (D), ±0.1 MM (±0.005 IN.) mm in.

3.1 2 3.4 2 3.8 8 4.1 7 4.5 4 4.9 5 5.3 3 5.7 2 6.2 2

0.12 3 0.13 5 0.15 3 0.16 4 0.17 9 0.19 5 0.21 0 0.22 5 0.24 5

1.4 7 1.5 7 1.6 3 1.7 3 2.0 3 2.1 3 2.2 8 2.5 4 2.8 4

0.05 8 0.06 2 0.06 4 0.06 8 0.08 0 0.08 4 0.09 0 0.10 0 0.11 2

6.8 6 7.6 2 8.3 8 9.1 4 9.9 0 10. 4 11. 2 11. 9 13. 5

0.27 0 0.30 0 0.33 0 0.36 0 0.39 0 0.41 0 0.44 0 0.47 0 0.53 0

4.9 8 5.4 6 5.9 7 6.3 0 6.9 6 7.2 6 7.7 5 8.2 6 9.2 7

0.19 6 0.21 5 0.23 5 0.24 8 0.27 4 0.28 6 0.30 5 0.32 5 0.36 5

2.3 9 2.7 7 3.1 8 3.5 8 3.9 6 3.9 6 4.7 5 4.7 5 4.7 5

0.09 4 0.10 9 0.12 5 0.14 1 0.15 6 0.15 6 0.18 7 0.18 7 0.18 7

1.9 1 2.0 6 2.1 6 2.3 1 2.6 4 2.8 2 3.0 5 3.3 5 3.7 6

0.07 5 0.08 1 0.08 5 0.09 1 0.10 4 0.11 1 0.12 0 0.13 2 0.14 6

5.6 1 6.3 5 6.8 6 7.5 4 8.3 3 8.5 9 9.0 9 9.3 5 10. 3

Source: Ref 1

RADIUS (S )

PUNCH RADIUS (R ), ±0.1 MM (±0.005 IN.)

HEIGHT (P), ±2%

0.22 1 0.25 0 0.27 0 0.29 7 0.32 8 0.33 8 0.35 8 0.36 8 0.40 6

RECESS RADIUS (R), R = S/3

RECESS HEIGHT (H), H = H/3

mm

in.

mm

in.

1.6 5 1.8 3 1.9 8 2.1 1 2.3 1 2.4 1 2.5 7 2.7 4 3.0 7

0.06 5 0.07 2 0.07 8 0.08 3 0.09 1 0.09 5 0.10 1 0.10 8 0.12 1

0.48 3 0.50 8 0.53 3 0.58 4 0.68 6 0.71 1 0.76 2 0.83 8 0.94 0

0.01 9 0.02 0 0.02 1 0.02 3 0.02 7 0.02 8 0.03 0 0.03 3 0.03 7

TABLE 2 PROCESS REQUIREMENTS FOR PROJECTION WELDING OF A RANGE OF HEAVY-GAGE LOW-CARBON STEELS USS GAGE NO.

THICKNESS (A)(B)

mm

PROJECTION SIZE

DIAMETER HEIGHT

in.

mm

in.

mm

in.

MINIMUM MINIMUM SPACING, CENTERLINE TO CENTERLINE mm in.

SCHEDULE A: WELDING NORMAL-SIZED WELDS 9 3.89 0.153 8.38 0.330 1.57 0.062 44.5 8 4.17 0.164 8.89 0.350 1.73 0.068 45.7 7 4.55 0.179 9.91 0.390 2.03 0.080 48.3 6 4.95 0.195 1.04 0.410 2.13 0.084 50.8 5 5.33 0.210 1.12 0.440 2.34 0.092 53.3 4 5.72 0.225 1.19 0.470 2.54 0.100 33.0 3 6.22 0.245 1.35 0.530 2.84 0.112 63.5 SCHEDULE B: WELDING SMALL-SIZED WELDS 9 3.89 0.153 6.86 0.270 1.47 0.058 40.6 8 4.17 0.164 7.37 0.290 1.57 0.062 41.9 7 4.55 0.179 7.87 0.310 1.70 0.067 43.2 6 4.95 0.195 8.38 0.330 1.83 0.072 44.5 5 5.33 0.210 8.89 0.350 1.96 0.077 45.7 4 5.72 0.225 9.40 0.370 2.08 0.082 48.3 3 6.22 0.245 9.91 0.390 2.24 0.088 53.3

MINIMUM CONTACT OVERLAP

ELECTRODE FORCE WELD

ELECTRODE FORCE FORGE

UPSLOPE TIME(D), CYCLES

WELDTIME(D), CYCLES

mm

in.

mn

lbf

mn

lbf

1.75 1.80 1.90 2.00 2.10 1.30 2.50

22.9 24.1 25.4 26.7 29.2 30.5 33.0

0.9 0.95 1.0 1.05 1.15 1.20 1.30

8.9 10.2 11.7 13.0 14.2 16.1 17.4

2000 2300 2630 2930 3180 3610 3900

17.8 20.5 23.4 26.1 28.3 32.1 34.7

4000 4600 5260 5860 6360 7220 7800

15 15 20 20 25 25 30

60 70 82 98 112 126 145

1.60 1.65 1.70 1.75 1.80 1.90 2.10

19.1 20.3 21.6 22.9 24.1 25.4 27.9

0.75 0.80 0.85 0.90 0.95 1.00 1.10

6.2 6.3 6.7 7.1 7.7 8.3 9.3

1400 1425 1500 1600 1730 1870 2100

12.5 12.7 13.4 14.2 15.4 16.6 18.7

2800 2850 3000 3200 3460 3740 4200

15 15 20 20 25 25 30

60 70 82 98 112 126 145

WELDING CURRENT(C), A

TENSILESHEAR STRENGTH(F)

mpa

ksi

15,400 16,100 17,400 18,800 20,200 21,500 23,300

52 56 66 78 86 103 119

7.5 8.1 9.5 11.3 12.5 15.0 17.3

11,100 11,800 12,800 13,900 14,900 16,000 17,300

35 38 45 53 59 72 83

5.1 5.5 6.5 7.7 8.5 10.4 12.0

Source: Ref 1

(A) LOW-CARBON STEEL, SAE 1005-1010, 290-380 MPA (42-55 KSI) ULTIMATE TENSILE STRENGTH. (B) SURFACE OF STEEL MAY BE OILED LIGHTLY BUT FREE FROM GREASE, SCALE, AND DIRT. (C) ON SINGLE FORCE WELDS USE ONLY WELD FORCE AS ELECTRODE FORCE. ELECTRODE FORCE CONTAINS NO FACTOR TO FURTHER FORM POORLY MADE PARTS. (D) BASED ON 60 HZ. (E) STARTING VALUES SHOWN ARE BASED ON EXPERIENCE OF RWMA MEMBER COMPANIES. (F) TENSILE-SHEAR STRENGTH PER PROTECTION DEPENDS ON THE JOINT DESIGN.

Slope control is recommended to prevent preflashing of the projection upon initiation of the welding current. Forge forces are also recommended, because weld porosity is often a concern with heavy-gage steel. Electrode or die wear may also be a concern, as it is in the spot welding of heavy-section steels. In such applications, pulsation welding schedules might be recommended. Intermediate-Gage Sheet Steels. The projection welding of steels ranging in gage from 0.56 to 3.43 mm (0.022 to 0.135 in.), using single-point projections with single-impulse welding schedules, is well established. Projection-stamping die designs are given in Table 3, whereas process requirements are given in Table 4. Lower forces and shorter welding times are required, when compared with the conventional spot welding of these gages.

TABLE 3 PROJECTION AND DIE GEOMETRIES FOR WELDING A RANGE OF INTERMEDIATE-GAGE STEELS TO MAKE SPHERICAL PROJECTIONS Data are for tool steel hardened to 50-52 HRC.

USS GAGE NO.

25-21 20-19 18-17 16-15 14 13 12 11 10

MATERIAL THICKNESS (T )

PROJECTION HEIGHT DIAMETER (H ), (D ), ±5% ±2%

mm 0.560.86 0.911.1 1.21.4 1.51.7 1.9 2.3 2.72 3.12 3.43

mm in. mm in. mm in. mm in. mm in. mm in. 0.64 0.025 2.3 0.090 9.53 0.375 0.79 0.031 1.93 0.076 2.29 0.090

in. 0.0220.034 0.0360.043 0.0490.054 0.0610.067 0.077 0.092 0.107 0.123 0.135

PUNCH DIAMETER (A)

POINT RADIUS (R ), ±0.05 MM (±0.002 IN.)

DIE HOLE DIAMETER (B ), ±0.1 MM (±0.005 IN.)

CHAMBER DIAMETER (D)

0.89 0.035 2.8

0.110 9.53 0.375 1.19 0.047 2.26 0.089 2.79 0.110

0.97 0.038 3.6

0.140 9.53 0.375 1.19 0.047 2.64 0.104 3.56 0.140

1.1

0.042 3.8

0.150 9.53 0.375 1.57 0.062 3.05 0.120 3.81 0.150

1.2 1.3 1.4 1.5 1.6

0.048 0.050 0.055 0.058 0.062

0.180 0.210 0.240 0.270 0.300

4.6 5.3 6.1 6.9 7.6

9.53 12.7 12.7 12.7 12.7

0.375 0.500 0.500 0.500 0.500

1.57 1.98 1.98 2.39 2.77

0.062 0.078 0.078 0.094 0.109

3.66 4.37 4.98 5.61 6.35

0.144 0.172 0.196 0.221 0.250

4.57 5.33 6.10 6.86 7.62

0.180 0.210 0.240 0.270 0.300

Source: Ref 1

TABLE 4 PROCESS REQUIREMENTS FOR PROJECTION WELDING OF A RANGE OF INTERMEDIATEGAGE LOW-CARBON STEELS Data are for SAE 1005-1010 steels with ultimate tensile strengths ranging from 290-380 MPa (42-55 ksi). Surface of steel may be oiled lightly but must be free from grease, scale, and dirt.

CONFIGURATION(A)

A B C D E F G H I J K CONFIGURATION(A)

USS GAGE NO.

25 23 21 19 18 16 14 13 12 11 10

MATERIAL THICKNESS

PROJECTION SIZE

DIAMETER HEIGHT

MINIMUM SPACING CENTERLINE TO CENTERLINE


mm 0.559 0.711 0.864 1.09 1.24 1.55 1.96 2.34 2.72 3.12 3.43

mm 2.29 2.29 2.79 2.79 3.56 3.81 4.57 5.33 6.10 6.86 7.62

mm 9.65 9.65 12.7 12.7 19.1 19.1 22.4 26.9 31.8 38.1 41.4

mm 6.35 6.35 9.65 9.65 12.7 12.7 12.7 15.7 19.1 20.6 22.4

WELD TIME(B), CYCLES

in. 0.022 0.028 0.034 0.043 0.049 0.061 0.077 0.092 0.107 0.123 0.135

in. 0.090 0.090 0.110 0.110 0.140 0.150 0.180 0.210 0.240 0.270 0.300

ELECTRODE FORCE(C)

mm 0.635 0.635 0.889 0.889 0.965 1.07 1.22 1.27 1.40 1.47 1.57

in. 0.025 0.025 0.035 0.035 0.038 0.042 0.048 0.050 0.055 0.058 0.062

WELDING CURRENT(D), A

in. 0.38 0.38 0.50 0.50 0.75 0.75 0.88 1.06 1.25 1.50 1.63

TENSILE-SHEAR STRENGTH(E)

kn lbf mpa ksi (F) WELDING SCHEDULE A (FOR SINGLE PROJECTION) A 3 0.67 150 4,400 2.6 0.37 B 3 0.87 195 5,500 3.4 0.50 C 3 1.1 240 6,600 4.8 0.70 D 5 1.5 330 8,000 7.3 1.06 E 8 1.8 400 8,800 9.0 1.30 F 10 2.4 550 10,300 12.4 1.80 G 14 3.6 800 11,850 16.7 2.43 H 16 4.5 1020 13,150 22.4 3.25 I 19 5.6 1250 14,100 26.5 3.85 J 22 6.7 1500 14,850 33.1 4.80 K 24 7.3 1650 15,300 37.9 5.50 WELDING SCHEDULE B (FOR 1-3 PROJECTIONS), EACH PROJECTION(G) A 6 0.67 150 3,850 2.3 0.33 B 6 0.67 150 4,450 3.0 0.43 C 6 0.67 150 5,100 3.7 0.53 D 10 0.93 210 6,000 6.1 0.88 E 16 1.2 270 6,500 7.6 1.10 F 20 1.6 365 7,650 11.0 1.58 G 28 2.4 530 8,850 14.8 2.15 H 32 3.0 680 9,750 19.3 2.80 I 38 3.7 830 10,600 23.8 3.45 J 45 4.5 1000 11,300 29.0 4.20 K 48 4.9 1100 11,850 33.4 4.85 WELDING SCHEDULE C (FOR [GES]3 PROJECTIONS), EACH PROJECTION(H)

in. 0.25 0.25 0.38 0.38 0.50 0.50 0.50 0.62 0.75 0.81 0.88

B C D E F G H I J K

8 11 15 19 25 34 42 50 60 66

0.45 0.56 0.71 0.98 1.5 2.1 2.7 3.3 4.0 4.5

100 125 160 220 330 470 610 740 900 1000

3300 3800 4300 4600 5400 6400 7200 8300 9200 9900

2.3 2.9 5.0 6.1 8.5 12.1 16.0 20.0 24.8 29.3

0.34 0.43 0.72 0.88 1.23 1.75 2.33 2.90 3.60 4.25

(A) CONFIGURATION DESIGNATIONS ARE FOR THE PURPOSES OF THIS TABLE ONLY. (B) BASED ON 60 HZ. (C) ELECTRODE FORCE CONTAINS NO FACTOR TO FURTHER FORM POORLY MADE PARTS. (D) STARTING VALUES SHOWN ARE BASED ON EXPERIENCE OF RWMA MEMBER COMPANIES. (E) TENSILE-SHEAR STRENGTH PER PROJECTION DEPENDS ON THE JOINT DESIGN. (F) SCHEDULE A IS USABLE FOR WELDING MORE THAN ONE PROJECTION IF CURRENT IS DECREASED, BUT EXCESSIVE WELD EXPULSION MAY RESULT AND POWER DEMAND WILL BE GREATER THAN THAT FOR SCHEDULE B OR C. (G) SCHEDULE B IS USABLE FOR WELDING MORE THAN THREE PROJECTIONS, BUT SOME WELD EXPULSION MAY RESULT, AND POWER DEMAND WILL BE GREATER THAN THAT FOR SCHEDULE C. (H) SCHEDULE C IS USABLE FOR WELDING LESS THAN THREE PROJECTIONS WITH WELD CURRENT INCREASED APPROXIMATELY 15% AND POSSIBLE OBJECTIONABLE FINAL SHEET SEPARATION. Thin-Gage Sheet. The projection welding of thin-gage steel sheet (75

1-3 >3

1 50

A parameter of particular interest to induction heating is the penetration depth of heating (also called the skin depth), and this is denoted by the symbol δ. The skin depth is defined as the position in the solid away from the surface at which the magnitude of the magnetic field falls to approximately 37% of its value at the part surface. The skin depth is given by: δ =

ρ πµ f

(EQ 2)

where ρ is the resistivity of the part, ρis the absolute magnetic permeability, and f is the frequency. The absolute permeability is equal to the product μrμ0, where μr is the relative permeability of the material (e.g., pure iron = 18,000 and nonmagnetic copper = 1) and μ0 is the permeability of "free space," which equals 1.257 × 10-6 V · s/m · A (a universal constant). Skin depth increases as the material resistivity increases and as the frequency decreases. For magnetic materials at temperatures below the Curie point, the value of μr is not a constant; it is dependent on the magnetic field intensity, H0. For example, the value of μr for a low-carbon 1044 steel is 734,000/ H 00.92 .


1. P. SIMPSON, INDUCTION HEATING, MCGRAW-HILL, 1960


1. P. SIMPSON, INDUCTION HEATING, MCGRAW-HILL, 1960 Induction Soldering Paul T. Vianco, Sandia National Laboratories

Safety Concerns The safety hazards accompanying the use of induction heating include electrical shock from contact with the coils as well as burns caused by the heated workpiece. A third hazard is the danger of steam explosion caused by arcing between the coil and the part. Typically, the heating coils are hollow copper tubing; they are cooled by the internal flow of water or a water-base coolant. The vapor cloud accompanying the spatter of flux during soldering can cause an electric arc that may puncture the copper tubing if the induction coil is close to the workpiece. Arcing can also be caused by movement of the coil windings through forces generated by the magnetic field. Water contacting the hot part produces an immediate, violent steam explosion. To prevent this occurrence, adequate clearance should be maintained between the coil and the part (typical minimum spacing is 6.4 to 13 mm, or 0.25 to 0.50 in.). Moreover, trial heating runs, in which the current is slowly increased to assess coil movement and part-to-coil shorting, should be performed prior to product assembly. Induction Soldering Paul T. Vianco, Sandia National Laboratories

References

1. P. SIMPSON, INDUCTION HEATING, MCGRAW-HILL, 1960 2. F. LOBKOWICZ AND A. MELISSINOS, PHYSICS FOR SCIENTISTS AND ENGINEERS, SAUNDERS CO., 1975, P 355 Wave Soldering Paul T. Vianco, Sandia National Laboratories

Introduction WAVE SOLDERING is one of the primary techniques for mass assembly of printed wiring boards involving through holes, surface mount devices, or a combination of these two technologies. A schematic of the wave soldering process is shown in Fig. 1. A solder "fountain" or "wave" is created by a pump located at the bottom of the solder pot; suitable baffles are mounted in the pot to direct the flow of solder into the desired configuration. The printed wiring board is placed onto a conveyor, which brings it into contact with the wave surface. Note from Fig. 1 that the circuit boards travel along the surface of the solder; molten alloy does not flow on top of the board. As the printed circuit board passes on the wave, the solder wets the surface-mount package leads, terminations, and exposed metal surfaces in the circuit board, and also fills plated through holes. This technique can produce several thousand solder joints in a matter of minutes.

FIG. 1 SCHEMATIC OF THE WAVE SOLDERING PROCESS. THE THREE IMPORTANT PROCESS CONTROL REGIONS ARE THE ENTRY (A), THE INTERIOR (B), AND THE PEEL-BACK REGION (C).

The implementation of wave soldering for surface-mount technology requires that the devices be glued to the substrate prior to wave soldering. Moreover, the entire surface-mount package must be able to withstand contact with the flux and temperatures of the molten solder without a loss of reliability. Small-outline transistors and discrete components (chip resistors, capacitors, etc.) are readily attached by wave soldering; larger packages (e.g., leaded or leadless chip carriers) may be damaged by exposure to the harsh environments (flux and molten solder) as well as be more prone to solder defects (Ref 1). The wave soldering technique encompasses a sequence of processes, all of which are typically contained in the same apparatus. First, the substrate receives a coating of flux. Flux application methods include: • •

•

A WAVE TECHNIQUE SIMILAR TO THAT SHOWN IN FIG. 1 COATING THE CIRCUIT BOARD BY A FLUX FOAM CREATED BY PASSING AIR OR NITROGEN THROUGH A FLUX BATH TO GENERATE THE FOAM (OR FROTH) ON THE BATH SURFACE DIRECTLY SPRAYING THE FLUX ONTO THE BOARD SURFACE

After flux is applied, the substrate is passed through a preheating stage. Warming the board promotes activation of the flux, accelerates the evaporation of volatiles from the flux (which can cause voids and solder balls by spattering upon contact with the hot solder bath), and reduces thermal shock to the substrate and devices when it passes onto the solder wave. Then, the circuit board contacts the solder wave for the formation of the joints. After passing the wave, the board cools through natural heat loss or, more quickly, by the use of forced air. Acknowledgement This work was performed at Sandia National Laboratories and supported by the U.S. Department of Energy under contract DE-AC04-76DP00789.

Reference

1. R. PRASAD, SURFACE MOUNT TECHNOLOGY: PRINCIPLES AND PRACTICES, VAN NOSTRANDRHEINHOLD, 1989, P 426-428

Wave Soldering Paul T. Vianco, Sandia National Laboratories

Design Considerations and Process Parameters The successful assembly of a printed circuit board by the wave soldering technique relies on the proper process parameters and the layout of the board for the respective technology (through-hole, surface-mount, or mixed). The hole or via size (diameter) must permit adequate capillary filling by the molten solder and form the joint or to interconnect within the short time period (5 to 7 s) that the board contacts the wave. Holes with small diameters (1.0 mm, or 0.039 in.) and gaps (>0.40 mm, or 0.016 in.) decrease the capillary driving force of the solder to fill the volume and create the joint; in some cases, the solder simply drains from the hole under its own weight. In addition to insufficient solder in holes and vias, other defects in wave-soldered products include bridges, icicles, and skips. Bridges are defects in which a quantity of solder forms a short across two conductors (leads, lands, or traces). A piece of solder that does not short two conductors, but rather hangs from one surface, is an icicle. Skips refer to the failure of solder in the wave to reach the metal surface and form the joint. Design Considerations Design considerations, such as the orientation of through-hole leads and surface-mount devices with respect to the direction of travel past the wave, as well as the proximity of packages to 0one another, can dramatically affect yields (Ref 1, 2). Figure 2 is a diagram showing the preferred orientation of through-hole and surface-mount packages.

FIG. 2 SCHEMATIC OF PREFERRED DEVICE ORIENTATION FOR WAVE-SOLDERED ASSEMBLIES. SOIC, SMALLOUT-LINE INTEGRATED CIRCUIT (SURFACE-MOUNT); DIP, DUAL-INLINE PACKAGE (THROUGH-HOLE)

Through-Hole Technology. In the case of through-hole assemblies, the leads on the side of a through-hole 16-pin dual-

in-line package (DIP) should be oriented perpendicular to the travel direction of the conveyor to prevent bridging between them. Moreover, the bridging of leaded packages can be limited by maintaining the lead extension beneath the circuit board at less than 1.0 mm (0.039 in.). Leads that are too close to each other also increase the chances of solder bridges. Solder skips are less frequently observed on through-hole products (assuming adequate wettability of the metal surfaces) because the device package body is on the top surface, and therefore does not shadow joints from the wave.

Surface-Mount Technology. The circumstances of solder defect function are different for surface-mount technology

because the entire package passes through the wave. Solder joint skips are more common because the package body can shadow nearby leads or terminations from contact with the wave, causing an inadequate supply of solder at the joint. Therefore, the joint areas of the package should be fully exposed to the wave, as suggested by Fig. 2. Packages with leads or terminations on all four sides experience an increased propensity for defects on the leads in the nonpreferred orientation. Also, the greater density of devices on surface-mount circuit boards increases the frequency of solder skips due to the closer proximity of the packages. Finally, the finer lead pitch used on surface-mount leaded and leadless (chip carrier) packages limits or prohibits their assembly on boards by wave soldering. Process Parameters The second factor critical to successful implantation of wave soldering of printed circuit boards is the process parameters. Process parameters can be divided into three groups: the fluxing operation, the solder wave properties, and the process schedule. The fluxing operation not only includes the particular parameters (wave or foam height, spraying pressure, etc.) needed

to ensure adequate flux supply to the board for solder wetting, but also the condition of the flux. In-line fluxing techniques (wave or foam) frequently use the same quantity of flux to process a large number of boards. Deterioration of the flux due to board contaminants, as well as loss of the flux vehicle, must be monitored to minimize defects in the final product. Specific gravity is used to indicate flux condition. Low-solids fluxes are gaining popularity as a "no-clean" alternative to rosin-base fluxes, including their use in wave soldering. The low-solids contents (typically less than 5%) make these fluxes difficult to foam, so spray or wave applications are preferred (Ref 3). Tighter control of the flux chemistry, reduced solder pot temperatures (by 10 to 15 °C, or 18 to 27 °F), shorter solder dwell times (by 1 to 2 s), and the use of inert or reducing atmospheres (Ref 4) can reduce the number of defects. Solder Wave Properties. The second process factor is the solder wave properties. The simple wave geometry (also

called the T-wave) is shown in Fig. 1. Three regions of the wave are identified. The entry region (on the left with the board moving to the right) should be sufficiently turbulent to supply solder to all of the exposed metal surfaces and should permit the escape of flux volatiles that may cause voids in the joint. In the case of surface-mount assemblies, fluid motion must be sufficiently active to force solder close to the packages for adequate wetting of terminations and leads, thereby avoiding solder skips. The interior region of the wave acts as the heat source and solder supply to promote wetting of the substrate features. The exit or peelback region is the point at which the board leaves the wave and is critical in preventing solder bridge and icicle defects on the circuit board. For example, defects are reduced by having a smooth wave that separates very slowly from the board. This effect is achieved by matching the solder flow velocity with the speed of the circuit board (Ref 5) as well as by making adjustments to the wave configuration. Hot air knives positioned after the wave have been used to remove solder bridges and icicles prior to solidification of the solder. Figure 3 shows several equipment modifications used to produce different wave geometries. The extended T-wave and lambda-wave modifications to the standard T-wave cause a smooth surface at the exit region to limit bridges, icicles, and the formation of excessive dross. The dual-wave configuration provides separate turbulent and smooth (laminar) waves to effect the desired solder joint qualities. Note that these modifications generally require that the circuit board laminate and devices withstand a longer period of contact with the molten solder.

FIG. 3 EQUIPMENT MODIFICATIONS USED TO PRODUCE DIFFERENT WAVE GEOMETRIES. THE NOMENCLATURE DESCRIBES THE PROFILE GEOMETRY OF THE WAVE.

Dross formation tends to be high in wave soldering due to the large surface areas of solder exposed to the atmosphere (which is further aggravated by the turbulent flow of the solder in the pot). Dross particles can become embedded in the solder joints as well as disturb the peel back of the solder upon exiting the board. The use of inert atmospheres, reduced solder temperatures, and oils to cover the solder surface limit the formation of dross. Solder temperatures are typically higher than those used in reflow processes. The higher temperatures are required to provide sufficient heat transfer to the substrate for adequate wetting as well as to reduce the solder surface tension to limit defects at the peel-back region. Solder pot temperatures are typically 50 to 80 °C (90 to 145 °F) above the liquidus temperature of the alloy; higher values are required for thicker substrates (e.g., multilayer or controlled-expansion substrates). The potential for damage to the laminate, including warpage and distortion as well as delamination of the conductor, increases with higher solder temperatures. Continuous processing with wave soldering equipment can cause a build-up of contaminants in the solder pot (e.g., nickel, copper, and gold from circuit board and lead finishes). See the article "Soldering in Electronic Applications" in this Volume for guidelines on acceptability limits for solder pot contamination. In addition to the traditional tin-lead alloys, wave soldering has been adapted to other solders, including the eutectic tinbismuth solder and the tin-lead-bismuth alloys (Ref 6). Process Schedule. The third factor in the wave soldering technique is the process schedule, which includes the preheat

zone temperature profile, the solder wave temperature, the conveyor tilt angle (α in Fig. 1, which affects the exit geometry of the board), and the conveyor speed. The conveyor speed determines the dwell time of the circuit board in the preheat zone and solder wave. Although specific settings depend on the type of equipment and the product being assembled, some general guidelines can be applied. First, the preheat stage should bring the board temperature to the range of 100 to 150 °C (210 to 300 °F). These temperatures provide suitable activation of the flux and lessen thermal shock to the circuit board (and devices) upon contact with the solder wave. The dwell time in the molten solder wave should be limited to 3 to 5 s. Dwell times at the molten solder temperature can reach 5 to 10 s for dual-wave systems. Time-temperature exposure limits for soldering processes are typically specified for devices and the laminate; manufacturer recommendations should

always be followed. Also, dwell times should be minimized to limit the leaching of conductor or lead finishes (e.g., copper, nickel, or gold) into the solder.


1. R. PRASAD, SURFACE MOUNT TECHNOLOGY: PRINCIPLES AND PRACTICES, VAN NOSTRANDRHEINHOLD, 1989, P 426-428 2. R. GENGLER AND J. HABIB, MINIMIZING DEFECTS IN THE MASS SOLDERING OF PRINTED WIRING ASSEMBLIES, WESTERN ELECTRIC ENGINEER, 1983 3. D. KOCKA, NO-CLEAN FLUXES ARE A VIABLE ALTERNATIVE TO CFC CLEANING,ELECT. PACK. PROD., JUNE 1990, P 95 4. P. FODOR AND P. LENSCH, COVER GAS SOLDERING LEAVES NOTHING TO CLEAN OFF PCB ASSEMBLY, ELECT. PACK. PROD., APRIL 1990, P 64 5. R. BOTHAM, C. LOWELL, AND J. STERRITT, WAVE SOLDERING MIXED TECHNOLOGY BOARDS, ELECT. PACK. PROD., NOV 1990, P 28 6. M. NYLEN, K. JOSEFSON, AND H. STEEN, USE OF TIN-LEAD-BISMUTH ALLOY AS A SUBSTITUTE FOR EUTECTIC TIN-LEAD IN WAVE SOLDERING, BRAZING AND SOLDERING, SPRING 1988, P 38 Wave Soldering Paul T. Vianco, Sandia National Laboratories

Defects Certain defect trends can be identified with particular process conditions (Ref 7). For example, voids and blowholes are more frequent at fast conveyor speeds and higher solder temperatures. These process conditions generate an excessively fast heating rate, which causes the rapid volatilization of the flux components and subsequent void formation. Rapid conveyor speed can be responsible for several other defects: •

•

•

AN INCREASED PROPENSITY FOR POOR FILLING OF HOLES AND VIAS CAUSED BY INSUFFICIENT DWELL TIME IN THE SOLDER WAVE. AS A RESULT, THE WETTING KINETICS MAY NOT FILL THE VOLUME (THIS IS FURTHER EXAGGERATED BY LOW SOLDER TEMPERATURES) THE PRESENCE OF SOLDER BRIDGES AND ICICLES DUE TO POOR MATCHING OF THE SUBSTRATE SPEED AND FLOW OF THE SOLDER WAVE AT THE PEEL-BACK REGION UPON BOARD EXIT CRACKING OF SURFACE-MOUNT CERAMIC CHIP CAPACITORS CAUSED BY THERMAL SHOCK

High solder pot temperatures and slow conveyor speeds (i.e., long dwell times) can create difficult-to-remove flux residues. Also, excessive warpage or distortion of the board results from inadequate support of the substrate by the conveyor system; this is particularly acute with thinner laminates, higher soldering temperatures, or longer dwell times (e.g., dual-wave systems). Thicker circuit boards (backplanes) may also warp due to the large thermal gradients through the thickness of the laminate due to inadequate preheating of the board. Clearly, the various process parameters can interact to produce certain defects. Therefore, optimization experiments and statistical process control are required to obtain and maintain acceptable product yields. The qualitative guidelines cited above provide a first measure of correlating defects to process conditions for subsequent improvement to the assembly procedure.


7. C. LEA, A SCIENTIFIC GUIDE TO SURFACE MOUNT TECHNOLOGY, ELECTROCHEMICAL PUBLISHERS, LTD., 1988, P 148-153 Wave Soldering Paul T. Vianco, Sandia National Laboratories

References

1. R. PRASAD, SURFACE MOUNT TECHNOLOGY: PRINCIPLES AND PRACTICES, VAN NOSTRANDRHEINHOLD, 1989, P 426-428 2. R. GENGLER AND J. HABIB, MINIMIZING DEFECTS IN THE MASS SOLDERING OF PRINTED WIRING ASSEMBLIES, WESTERN ELECTRIC ENGINEER, 1983 3. D. KOCKA, NO-CLEAN FLUXES ARE A VIABLE ALTERNATIVE TO CFC CLEANING,ELECT. PACK. PROD., JUNE 1990, P 95 4. P. FODOR AND P. LENSCH, COVER GAS SOLDERING LEAVES NOTHING TO CLEAN OFF PCB ASSEMBLY, ELECT. PACK. PROD., APRIL 1990, P 64 5. R. BOTHAM, C. LOWELL, AND J. STERRITT, WAVE SOLDERING MIXED TECHNOLOGY BOARDS, ELECT. PACK. PROD., NOV 1990, P 28 6. M. NYLEN, K. JOSEFSON, AND H. STEEN, USE OF TIN-LEAD-BISMUTH ALLOY AS A SUBSTITUTE FOR EUTECTIC TIN-LEAD IN WAVE SOLDERING, BRAZING AND SOLDERING, SPRING 1988, P 38 7. C. LEA, A SCIENTIFIC GUIDE TO SURFACE MOUNT TECHNOLOGY, ELECTROCHEMICAL PUBLISHERS, LTD., 1988, P 148-153 Vapor-Phase Soldering Dale L. Linman, Centech Corporation

Introduction VAPOR-PHASE SOLDERING is a process of condensation heating, in which a product prepared for soldering is passed through or into a layer of saturated vapor. The vapor condenses on the relatively cool part, which transfers the latent heat of vaporization to the part and heats it rapidly and uniformly. The process reaches thermal equilibrium in a few seconds at a temperature that is only limited by the boiling point of the fluid used in the process. Typically, newer equipment utilizes a single perfluorocarbon fluid vapor layer, which does not contain chlorine or bromine and is environmentally safe. This fluid is also chemically inert, thermally stable, and has very low solvent action. Although the process is used in a variety of heating applications, such as curing epoxies and polymers or stress relieving various platings, it is primarily useful when soldering surface-mounted components to various substrate materials. Because of the precise temperature control that this process offers (by virtue of a fixed boiling point), multilevel soldering is possible using up to three solder compositions with three different melting temperatures. Fluid temperatures that range from 100 to 265 °C (212 to 510 °F) are available. The most common fluid used boils at 215 °C (420 °F) and is used for 63-37 type solder, which melts at 183 °C (360 °F). A 30 °C (55 °F) differential is typical, to enable the heating process to proceed rapidly. Temperature precision also prevents the assembly from overheating. Vapor-Phase Soldering

Dale L. Linman, Centech Corporation

Typical Materials The typical materials that are joined by this process are solder-coated substrates made from various ceramic materials and G10 or FR4 epoxy laminates mated with surface-mounted components. Solder paste is typically applied to the substrate by screening, stenciling, or using X-Y dot placement. Components are then placed in the correct position in the solder paste, which temporarily holds them until they are soldered. The assembly is then passed through the vapor-phase setup, where the solder is melted, joining the components to the substrate. Typical surface-mount lead density is 1.25 mm (0.050 in.) on centers. Fine-pitch components as small as 0.38 mm (0.015 in.) on centers can also be processed. There is no inherent lead-pitch limitation in the vapor-phase process, because it is essentially an oxygen-free soldering process. Because of its uniformity, rapid heating, precise temperature control, and oxygen-free environment, the process is preferred for high-value assemblies. Also, because of the thermal capacity, ease of profiling, and forgiving nature of the process, it is also used when high throughput is required from limited floor space or when a large number of different types of boards need to be run in a given time period. Equipment is normally configured for either batch or in-line operations (Fig. 1 and 2). The majority of machines available are of single-vapor design (Fig. 3). Earlier models had vertical access and used less-expensive fluid as a cover layer. These machines also lacked a preheating capability, which was found to be necessary in order to reduce the temperature differential at the point of solder melting. This measure prevents wicking and other undesirable characteristics that can develop when one part of a solder joint heats faster than the other. Both a preheating capability, typically provided by infrared panels, and a cooldown stage have been incorporated in newer models. The cooldown stage brings the solder joint through the liquidus phase more quickly (Fig. 4). This procedure reduces the time that the joint can form undesirable intermetallic layers and simultaneously improves the grain structure of the joint, thus improving its capability to withstand thermal cycling.

FIG. 1 PELLETIZED VAPOR-PHASE BATCH SYSTEM

FIG. 2 THIRD-GENERATION VAPOR-PHASE REFLOW SYSTEM

FIG. 3 SINGLE-VAPOR BATCH-TYPE EQUIPMENT

FIG. 4 THERMAL PROFILE

The stability of the vapor-phase process and its ease of profiling means that machine operators require little training. If the overall process is established correctly, then there are few elements that require adjustments. A smooth, vibration-free conveying mechanism enables defect-free soldering and assembly. The proper setup procedure is to attach a thermocouple to a strategic location on the printed board, such as a sensitive component or a solder joint. Then, the board is processed to reach a preheat temperature of between 125 and 150 °C (255 and 300 °F) in 50 to 60 s. A desirable heating rate is approximately a 2 °C/s (4 °F/s) rise to within 100 °C (180 °F) of the reflow temperature. Major parameters are dwell time (batch machines), conveyor speed (in-line machines), and preheat panel temperature. Reflow time is established by observing the solder melt through the window in the chamber. In the batch machine, the dwell times can be set from those obtained and the process replicated. With in-line machines, the reflow time is established by the conveyor speed and the preheat temperature is adjusted with the panel temperature. Similar board assemblies can be processed using very similar profiles, which minimizes lost machine time due to the complex repetitive profiling that is common with other reflow methods. Because vapor-phase soldering depends on the series of processes that precede it, every parameter leading to the actual soldering process must be carefully controlled to ensure a high yield. For example, thickness, viscosity, and location accuracy are just some of the parameters involved in solder paste application. Pad design and board layout are equally important to the process, as is the accurate placement of components. With proper controls, it becomes possible to solder with defect rates under 50 ppm using the vapor-phase process. Information on guidelines for parameter control can be obtained by contacting the Institute for Interconnecting and Packaging Electronic Circuits.

Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Introduction THE SELECTION of materials for welded construction applications involves a number of considerations, including design codes and specifications where they exist. Mobile structures, such as automobiles and aerospace vehicles, have quite different materials requirements for weight, durability, and safety than stationary structures, such as buildings and bridges, which are built to last for many years. In every design situation, economics--choosing the correct material for the life cycle of the part and its cost of fabrication--are of great importance. Design codes or experience frequently offers an adequate basis for material selection; for new or specialized applications, however, the engineer encounters problems of an unusual nature and thus must rely on basic properties of the material, such as strength, corrosion or erosion resistance, ductility, and toughness. The properties of the various metallurgical structures associated with the thermal cycles encountered in the welding operation must also be included in the design process. The various subsections in this article offer guidance for material selection applications involving bridges and buildings, pressure vessels and piping, shipbuilding and offshore structures, aerospace systems, machinery and equipment, automobiles, railroad systems, and sheet metal. Material properties and welding processes that may be significant in meeting design goals are described.

Note

* R. DAVID THOMAS, JR., CHAIRMAN, R.D. THOMAS & COMPANY; BRUNO L. ALIA; WILLIAM R. APBLETT. AMET ENGINEERING; ROBERT G. BARTIFAY, ALUMINUM COMPANY OF AMERICA; STEPHEN A. COUGHLIN, ACF INDUSTRIES, INC.; GREGORY MELEKIAN, GENERAL MOTORS CORPORATION; ANTHONY R. MELLINI, ST., MELLINI AND ASSOCIATES, INC.; LARRY PERKINS, WRIGHT LABORATORY; JAMES E. ROTH, JAMES E. ROTH, INC.; WILLIAM J. RUPRECHT, GENERAL ELECTRIC COMPANY; BERNARD E. SCHALTENBRAND, ALUMINUM COMPANY OF AMERICA; ROBERT E. SOMERS, SOMERS CONSULTANTS; ROGER K. STEELE, AAR TECHNICAL CENTER Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Overview

Principles of Material Selection Structural Integrity. Before a structure is initially placed in service, the properties of its materials of construction must be adequate to ensure that the structure will perform as intended. Its components must be strong enough to bear the loads that will be imposed. It must be capable of meeting various inspection and qualification tests required by the end user to measure base material and weld soundness, ductility, toughness, and strength. Length of Service. The service environment is often the largest influence on material selection. Materials of

construction rarely last indefinitely. One of the principal selection considerations is the intended length of service for the structure. The designer needs to choose materials that will meet the conditions imposed initially and that can adapt to possible changes in conditions during the life of the structure. The length of service is often dictated by obsolescence. Military hardware fabricated during wartime is an example of structures with relatively high obsolescence and short service life. Rocket cases for the propulsion of space vehicles need

to last one launching. At the other end of the spectrum, hydroelectric projects have expected lifetimes of a century or longer. Materials or components rarely reach the end of their useful lives at the same time. Although a hydroelectric project is designed to last 100 years, some components may reach the end of their expected lives in less time, either by obsolescence due to improved technology or by unavoidable wear. At the outset of the material selection process, the length-of-service criterion must take into consideration the following: inspectability for potential failure, ease of maintenance of areas subject to deterioration, and ease of replacement of obsolete or life-exhausted components. Joining Process. The suitability of the materials for the fabrication process to be employed must be considered early in

the selection process. Weldability is the property of a material that dictates its ease of joining. The term most often applies to various types of steels joined by arc welding processes. This property often determines whether preheat is needed for successful welding and whether some type of postweld heat treatment is required to restore properties degraded by the welding operation. Although weldability usually relates to the joining of relatively thick plates, the choice of materials for sheet metal applications must also take the welding process into account. The thinner the metal, the more difficult it is to arc weld. Brazing, soldering, or resistance welding is often used to join thin sheet materials. Resistance welding is frequently employed in the assembly of automobile bodies, for example; however, this process may not perform as reliably on coated steels as it does on uncoated steels. Design Codes. Material selection is often dictated by the codes by which the design is governed, but there is still

frequently a choice of materials. For example, pressure vessel codes offer a choice of steels, depending on the service temperature. Structural codes for buildings and bridges allow the use of both plain carbon and low-alloy steels. However, the codes tend to narrow the choice of materials. Economics. An overriding selection criterion is often cost. In the broadest sense, this can be determined by the price of

the material. However, a designer must assess not only the unit cost of the selected material, but also the costs imposed by fabrication. Compared with unalloyed steel, a low-alloy steel may allow a structure to be built using 25% less steel at a nominal 15% increase in unit price, but the cost of fabrication, such as the need for preheat and postweld heat treatment, may outweigh the material saving. A more costly steel that does not require painting or other protection from atmospheric corrosion may be a more economical choice, if not for initial construction, then for the life of the structure. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Service Conditions The service environment is a principal determinant in the selection of materials for welded construction applications. The physical, mechanical, and corrosion properties of a material must be considered (Table 1).

TABLE 1 PRINCIPAL ATTRIBUTES OF A MATERIAL

PHYSICAL PROPERTIES • • • • •

DENSITY MELTING POINT THERMAL EXPANSION ELECTRICAL CONDUCTIVITY FERROMAGNETISM

MECHANICAL PROPERTIES • • • • • • • •

TENSILE STRENGTH IMPACT STRENGTH FATIGUE STRENGTH CREEP STRENGTH DUCTILITY HARDNESS FRACTURE TOUGHNESS ELASTIC LIMIT

CORROSION PROPERTIES • • • • • •

GENERAL CORROSION PITTING ATTACK STRESS CORROSION EROSION CAVITATION OXIDATION

Ambient Temperatures and Pressures. Most welded structures are designed to work in normal atmospheric

conditions. The properties of common materials are well known for such conditions, and ample data are available to assist the designer in material selection. Elevated Temperatures. The chemical industry has many processes that operate at temperatures of 1000 °C (1830 °F)

and higher. Fossil fuel electric power generation makes use of steam temperatures from 500 to 600 °C (930 to 1110 °F). At such temperatures, most steels will exhibit creep--that is, a gradual deformation in the direction of the principal stresses--and thus the creep strength of a material becomes the controlling property. Creep strength determines the unit stress allowed to minimize the rate of creep and hence the permanent distortion and ultimate failure of a component. Materials operating at high temperatures are frequently exposed to corrosive, erosive, or oxidizing environments, which may determine the life span of a structure. The resistance of proposed materials of construction to these types of environments should be examined during the selection process. In processes where heat transfer occurs, the thermal conductivity of a material becomes an important consideration. Allowance must be made in the design for the expansion and contraction of components during heating and cooling-phenomena related to the coefficient of thermal expansion of a material. Low Temperatures. Most ferritic metals become brittle at low temperatures, losing their ability to deform plastically to

adjust to localized stresses and to accept a sudden shock. Thus, materials for use at low temperatures must have good toughness, a property usually measured by the Charpy V-notch impact test. For low-temperature applications, as for elevated-temperature service, the physical properties of thermal conductivity and coefficient of expansion become important. High Pressures. Structures intended to contain high pressures are usually weldments and are designed according to

American Society of Mechanical Engineers (ASME) codes for pressure vessels and piping. Strength properties are of primary importance. The codes specify the unit stress allowed for pressure-containing components. In cylindrical vessels, the maximum stress is in the tangential (i.e., hoop) direction, and the axial stress is half the hoop stress. In its simplest form, hoop stress is determined by: S=

PD 2T

where S is unit stress, P is pressure, D is the diameter of the cylinder, and T is wall thickness.

The tensile strength of a material determines the allowable stress and is specified in the ASME codes. The codes for steels allow a stress in excess of two-thirds of the minimum yield strength but less than 90% of the minimum yield strength. The minimum wall thickness, Tm, is then determined by: Tm =

PD +A 2 S + Pv

where the extra thickness allowance, A, provides for anticipated corrosion or wear, and the coefficient y is specified by the code and varies with the temperature of operation and the wall thickness relative to the diameter. Vacuum Conditions. Materials that are selected for very-low-pressure applications are primarily concerned with the

tightness of the joints to prevent leakage from the exterior higher (or ambient) pressures into the vacuum chamber. For welded containers, the properties of the weld metal are critical to ensure lack of defects, such as porosity, cracks, and microfissures. Vacuum conditions exist in outer space. In such applications, the vacuum condition is on the exterior of the vessel, and design for pressure containment is similar to that described for high pressures. Material selection also requires consideration of dynamic loads experienced during launches (hence, toughness), density for light weight (as in aircraft and aerospace applications), and corrosive and erosive effects, especially for reentry into the atmosphere. Underwater Applications. A wide variety of material properties are involved in the application of weldments under

water. Corrosion resistance is frequently a major concern, especially in seawater. Ships and offshore drilling platforms are notable examples. For ships and for piping operating in biologically active waters, microbiologic organisms pose an unusual problem in material selection. Barnacles that attach themselves to the hulls of ships greatly increase the resistance to propulsion. The use of copper cladding or paints containing copper-bearing compounds is one means of discouraging barnacle formation. In turbid waters, piping is sometimes attacked by microorganisms that form surface pits, causing cracks in areas of local tension stresses. Because high tensile stresses frequently occur in or adjacent to welds, failures of weldments in these conditions can be attributed to these organisms, even in materials that otherwise are considered corrosion resistant to such environments. Corrosion, Erosion, and Oxidation. Corrosion generally refers to surface attack by aqueous solutions or organic

chemicals in either the liquid or vapor phase. Erosion implies the mechanical removal of metal from the surface, sometimes by high-pressure impingement of a fluid, in which case it is termed cavitation. Oxidation commonly occurs in metals at elevated temperatures and results in the removal of metal in the form of oxides. Data on the resistance of materials to deterioration in many of these environments are frequently available from handbooks and material suppliers. In the case of weldments, the performance depends largely on experience gained in a variety of applications. The weld metal has a structure distinctly different from that of the parent metal, and sometimes a different composition as well. An electrochemical potential may exist between the weld and parent metal structures, causing attack on one or the other (frequently the weld metal because of its lower surface area). Corrosive attack in weldments often is attributed to stress-corrosion cracking (SCC). In the presence of chloride ions (even at levels of parts per million), high tensile stress zones in or adjacent to welds can cause failure in materials whose corrosion life in the absence of stress is more than adequate for a particular application. The welds in components carrying fluids at high velocity may erode during service. Under such conditions, hardness is the critical property. Localized areas of components, such as valve seats and turbine blades, may be protected from erosion by applying a hard weld metal, known as hard surfacing. For more information, see the article "Hardfacing, Weld Cladding, and Dissimilar Metal Joining" in this Volume. In high-temperature service, allowance must be made for the thinning of metal sections due to progressive oxidation. In some applications the initial oxide layer protects the underlying metal, and progress is arrested. Adjacent to welds, oxidation can form preferentially in the highly stressed zones adjacent to welds and progress through the wall, causing weakness and possible failure.

Fatigue. In weldments subjected to fluctuating stresses, the fatigue strength of a metal may be the primary consideration

in material selection and design loads. In the absence of flaws, the fatigue strength of most metals used for dynamically loaded structures is considerably higher than that of the welds. Designs for such structures must take into consideration both service experience and laboratory experiments dealing with various types of welds. Weld metal cracks (even fine microfissures), lack of penetration or lack of fusion at the interface, and geometric discontinuities on the surface (such as overlaps and undercuts) are the major causes of poor performance in fatigue tests of welds. In the absence of such flaws, the fatigue properties of the weld metal itself or of the weld heat-affected zone (HAZ) rarely need to be considered, even though these zones may be inferior compared with the wrought parent metals. Design allowances for fatigue of weldments primarily take into account potentially damaging flaws. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Effect of Weld Thermal Cycles Thermal Excursions During Welding. The thermal cycles encountered during the welding operation must be considered when selecting material for weldments. To make a weld requires that temperatures in the fusion zone reach above the melting point of the metal. Zones adjacent to the fusion region reach varying peak temperatures, depending on their distance from the fusion zone. Each of these thermal cycles causes metallurgical changes that affect the properties of the metal.

In the weld metal itself, the dendritic structure formed during solidification in a single thermal cycle is similar to that of a casting. In multipass welding, which is common when arc welding thicknesses greater than 6.4 mm (

1 in.), the weld 4

metal will undergo several thermal cycles, a process that refines the dendritic structure and alters the weld properties. Similarly, in the HAZs immediately adjacent to the weld, one or more heat cycles will be experienced, causing changes in properties. When selecting materials for welded construction, the properties of these HAZs must be considered. Cooling from the peak temperatures causes residual stresses, most often of a magnitude matching the yield strength of the material. Residual stresses may add to the service stresses. In cases where residual stresses may influence the integrity of the structure during its life expectancy, the designer may specify a postweld heat treatment. Postweld Heat Treatments. Many weldments are placed in service in the as-welded condition. Under certain

circumstances, postweld heat treatments may be required, thereby changing the properties of the selected metal. Stress relieving is a thermal cycle that takes a weldment up to a high temperature where its yield strength is substantially reduced, allowing localized plastic deformation of the highly stressed regions. In ferritic steels, most stresses are relieved by heating to a temperature of less than 720 °C (1330 °F) for carbon steels (note that alloy steel values are higher). This is below the temperature transformation of ferrite to austenite, known as the critical temperature. This softens the structures that have been hardened by fast cooling from above the austenite-forming temperatures. Such subcritical heat treatments are said to "temper" the steel. To cause the structures to become more similar to that of the parent metal, a supercritical heat treatment is sometimes specified. The properties of such welds are then likely to be comparable to those of the parent metal, provided that the compositions are similar. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Bridges and Buildings

This section deals with large welded structures built of plain carbon steels--the tonnage items in the construction industry. These structures generally conform to established codes. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Material Selection The four code-writing organizations listed in Table 2 specify the chemical compositions, tensile properties, and other pertinent characteristics of steels used for large welded structures. Most employ the steel specifications of ASTM.

TABLE 2 CODES FOR STEEL STRUCTURES

ORGANIZATION

TITLE

AMERICAN WELDING SOCIETY (AWS)

"STRUCTURAL WELDING CODE-STEEL" "BRIDGE WELDING CODE" "DESIGN, FABRICATION, AND ERECTION OF STRUCTURAL STEEL FOR BUILDINGS" "STANDARD SPECIFICATION FOR WELDING STRUCTURAL STEEL HIGHWAY BRIDGES" "MANUAL FOR RAILWAY ENGINEERING"

AMERICAN INSTITUTE OF STEEL CONSTRUCTION (AISC) AMERICAN ASSOCIATION OF STATE HIGHWAY AND TRANSPORTATION OFFICIALS (AASHTO) AMERICAN RAILWAY ENGINEERS ASSOCIATION (AREA)

SPECIFICATION NO. D1.1 D1.5 ...

...

...

The "Structural Welding Code--Steel" (AWS D1.1) considers three classes of structures: • • •

STATICALLY LOADED (BUILDINGS) DYNAMICALLY LOADED (BRIDGES) TUBULAR (OFFSHORE AND ARCHITECTURAL)

Fifteen ASTM specifications are applicable to buildings, 13 to bridges, and 34 to tubular structures, including four of the American Petroleum Institute (API) and two of the American Bureau of Shipping (ABS). The six most commonly used ASTM specifications, applicable to all three classes, are:

NO. A 36 A 500 A 501 A 514 A 572

TITLE "SPECIFICATION FOR STRUCTURAL STEEL" "SPECIFICATION FOR COLD-FORMED WELDED AND SEAMLESS CARBON STEEL STRUCTURAL TUBING" "SPECIFICATION FOR HOT-FORMED WELDED AND SEAMLESS CARBON STEEL STRUCTURAL TUBING" "SPECIFICATION FOR HIGH-YIELD STRENGTH, QUENCHED AND TEMPERED ALLOY STEEL PLATE, SUITABLE FOR WELDING" "SPECIFICATION FOR HIGH-STRENGTH LOW-ALLOY COLUMBIUM-VANADIUM STEELS OF STRUCTURAL QUALITY"

A 588 A 852

"SPECIFICATION FOR HIGH-STRENGTH LOW-ALLOY STRUCTURAL STEEL WITH 50 KSI (345 MPA) MINIMUM YIELD POINT TO 4 IN. THICK" "STANDARD SPECIFICATION FOR QUENCHED AND TEMPERED LOW-ALLOY STRUCTURAL STEEL PLATE WITH 70 KSI YIELD STRENGTH TO 4 IN. THICK"

These specifications vary considerably in properties important to notch sensitivity (toughness) and weldability. The summary in Table 3 provides a guide as to what to look for in these and other specifications referred to in the structural codes. The supplementary requirements in some of these specifications are useful to the designer and fabricator when evaluating the need for additional acceptance tests for unusual applications.

TABLE 3 PROPERTIES RELATING TO SELECTION FOR SERVICE OF ASTM STEELS APPROVED IN AWS D1.1-92 Includes only the seven steels approved for all three AWS structural classes: (1) statically loaded, (2) dynamically loaded, and (3) tubular designs.

ASTM MAXIMUM NO. CARBON CONTENT, % A 36 0.25-0.29 (FOR PLATE THICKER THAN 4 IN., OR 100 MM)

TENSILE TEST REQUIREMENTS(A)

DEOXIDATION

400-550 MPA (55-80 NO RIMMED OR KSI); 250 MPA (36 KSI) CAPPED STEEL MIN YP FOR PLATE THICKER THAN 13 MM (

A 500

0.23-0.30

A 501

0.26-0.30

A 514

0.10-0.21 (15 PROPRIETARY COMPOSITIONS)

A 516

0.18-0.27

A 572

0.21-0.26

A 588

0.15-0.20 (WEATHERING STEELS)

(A) (B)

1 2

SUPPLEMENTARY REQUIREMENTS CHARPY V-NOTCH TESTING; SILICON KILLED, FINE-GRAIN PRACTICE

IN.)

MIN TENSILE, THREE GRADES; 230, 290, 320 MPA (33, 42, 46 KSI) MIN YP 400 MPA (58 KSI) MIN TENSILE; 250 MPA (36 KSI) MIN YP 760-900 MPA (110-130 KSI) TO 65 MM (2.5 IN.) THICK; 690 MPA (100 KSI) MIN YP PLATES THICKER THAN 38 MM (1.5 IN.) NORMALIZED

...

...

...

...

FULLY KILLED, FINE-GRAIN PRACTICE

CLOSER THAN STANDARD FLATNESS

KILLED, FINEGRAIN PRACTICE

MIN TENSILE, FOUR GRADES; 290-450 MPA (42-65 KSI) MIN YP MIN TENSILE

...

VACUUM TREATMENT; CHARPY V-NOTCH TEST; DROP-WEIGHT TEST(B) MAX TENSILE

YP, YIELD POINT. SEE ASTM E 208.

FINE-GRAIN PRACTICE

MAX TENSILE; CHARPY V-NOTCH TEST; DROP-WEIGHT TEST(B); ULTRASONIC TEST

In general, the lower the carbon content and the finer the grain size of a steel, the better its notch toughness and weldability. The methods for achieving superior properties add to the unit price of the steel, so a designer striving for a low-cost structure tries to reduce steel cost, which is usually the most expensive single item in a job. This is especially true of static structures.


Environment Low Temperatures. In the design of bridges, certain components merit special-toughness steel. Such components, termed fracture critical tension members (FCMs), would allow the bridge to collapse should they fail. The steel selection criteria are based on the minimum Charpy V-notch energy at a given temperature. For bridges financed in part by the Federal Highway Administration (FHA), the minimum winter temperature is the criterion used for the required Charpy Vnotch energy at that temperature. A bridge located in the northern United States requires FCMs made with steel that exhibits the required minimum Charpy V-notch energy at a lower temperature than the same bridge constructed in the southern U.S.

The same FCM concept also applies for railroad bridges constructed under the codes established by AREA. The Charpy V-notch requirements for the three temperature zones defined by AREA codes are shown in Table 4. The required temperatures for measuring the Charpy V-notch energy are 21 °C (70 °F) for zone 1, 4 °C (40 °F) for zone 2, and -23 °C (-10 °F) for zone 3. Higher minimum Charpy V-notch values are required for weathering steels in heavier thicknesses. These values have been selected on the basis of extensive studies of steel structures, especially those of the Liberty ship failures during World War II.

TABLE 4 CHARPY IMPACT TEST REQUIREMENTS FOR WELDED STRUCTURAL STEEL: FRACTURE CRITICAL MEMBERS

ASTM

THICKNESS

A 36 OR A 709, GRADE 36

≤ 38

MM ( ≤ 1.5 IN.)

>38 TO ≤ 100 MM (>1.5 TO ≤ 4 IN.) A 572, GRADE 50

≤ 38

MM ( ≤ 1.5 IN.)

A 709, GRADE 50

>38 TO ≤ 50 MM (> 1.5 TO ≤ 2 IN.)

A 588 A 709, GRADE 50 W MINIMUM SERVICE TEMPERATURE

>38 TO ≤ 100 MM (>1.5 TO ≤ 4 IN.)

MINIMUM AVERAGE ENERGY AND TEST TEMPERATURE ZONE 1 ZONE 2 ZONE 3 34 J (25 FT · 34 J (25 FT · 34 J (25 FT · LBF) AT 21 °C LBF) AT -12 °C LBF) AT 4 °C (40 (70 °F) (10 °F) °F) 34 J (25 FT · 34 J (25 FT · 34 J (25 FT · LBF) AT -1 °C LBF) AT -7 °C LBF) AT -23 °C (30 °F) (20 °F) (-10 °F) 34 J (25 FT · 34 J (25 FT · 34 J (25 FT · LBF) AT 21 °C LBF) AT -12 °C LBF) AT 4 °C (40 (70 °F) (10 °F) °F) 34 J (25 FT · 34 J (25 FT · 34 J (25 FT · LBF) AT -1 °C LBF) AT -7 °C LBF) AT -23 °C (30 °F) (20 °F) (-10 °F) 41 J (30 FT · 41 J (30 FT · 41 J (30 FT · LBF) AT -1 °C LBF) AT -7 °C LBF) AT -23 °C (30 °F) (20 °F) (-10 °F) -18 °C (0 °F) -34 °C (-30 °F) -51 °C (-60 °F)

Source: Adapted from "Manual for Railway Engineering," Table 1.14.7. American Railway Engineers Association Earthquakes. Structures for locations subject to earthquakes should be designed to withstand the impact of such events.

This is a constantly evolving technology. For up-to-date information, contact the Earthquake Engineering Research Center (EERC) at the University of California at Berkeley. Atmospheric Corrosion. Steel structures that are to remain exposed to the atmosphere without painting or other surface

protection can be constructed of a steel that forms an adherent oxide. Such steel is known as weathering steel and is covered by ASTM specification A 588. Corrosion resistance is provided by additions of approximately 0.35% Cu in most grades and/or small amounts of chromium or nickel in other grades. Weathering steels are suitable for resisting corrosion in air, but are not significantly better than other structural steels under water. In highway bridges where salt is used to control icing conditions, progressive corrosion is encountered, especially at expansion joints. A matching filler metal should be used to weld these steels, such as those with the designation "W" in the AWS specifications (e.g., E7018W in A5.5 or EW in A5.23). Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Fitness for Service No structure is entirely free from flaws. Fitness for service is a term applied to an evaluation of the flaws found on inspection. The codes define the allowance for defects considered acceptable for the intended service. Workmanship. Prior to assembly, the base metals are inspected for evidence of surface flaws that could impair the integrity of the structure. Before welding, the prepared edges are examined for discontinuities longer than 25 mm (1 in.). Such flaws must be removed and weld repaired before final assembly.

Distortion and shrinkage result from the welding operation. Permissible variations in the dimensions, straightness, and camber of structural members are defined in the codes. Peening of intermediate layers of multi-pass welds can be used to control distortion, but should be avoided on the weld surfaces. Acceptable weld profiles are also defined as part of the workmanship standards of the codes for buildings and bridges. Unacceptable profiles, cracks, undercuts, overlaps, and other deficiencies are generally repaired by grinding the objectionable flaw and, if necessary, rewelding. Inspection. Visual inspection is relied on for assessing workmanship criteria. In the construction of large structures, the

general contract defines the extent of nondestructive testing of welds that may be required. Methods of nondestructive testing include radiography, ultrasonic testing, magnetic-particle inspection, and dye-penetrant inspection. For more information, see the article "Inspection of Welded Joints" in this Volume. Radiography and ultrasonic testing allow examination of the full cross section of a weld. Flaws larger than 1% of the thickness, including cracks, lack of fusion, inclusions, and porosity can be detected by these techniques. X-ray analysis will detect only cracks substantially parallel to the impinging x-ray beam. Ultrasonic pulses can be introduced at an angle of 30 or 45° from the perpendicular to the plate surface, which permits a more complete inspection coverage of a weld. Magnetic-particle and dye-penetrant methods are limited to the detection of defects on or near the surface. Dye penetrants can detect surface flaws only, whereas magnetic-particle testing can detect flaws slightly below the surface. Flaw Assessment. The integrity and performance of welded structures depend heavily on the assessment of the

significance of flaws identified during inspection. The discipline of fracture mechanics allows judgments to be made as to the likelihood of a flaw to cause a failure. For static structures, the acceptance of a flaw of a given size that is found during inspection makes use of the material property evaluated by the crack-opening displacement (COD) test. Acceptance is based on the principle illustrated in Fig. 1. An example of the application of the COD test to flaw assessment is shown in Fig. 2.

FIG. 1 SCHEMATIC OF AN ALLOWABLE FLAW-SIZE CURVE

FIG. 2 ALLOWABLE FLAW SIZES DETERMINED FOR GIRTH WELDS IN PIPELINES CONSTRUCTED FROM API X70 STEEL, ASSUMING A COD VARYING FROM 0.1 TO 0.25 MM (0.004 TO 0.010 IN.), A FLOW STRESS OF 520 MPA (75 SKI), AND AN APPLIED STRESS OF 435 MPA (63 KSI). SOURCE: REF 1

Tolerance for Flaws. The inherent material properties of steels and welds that allow them to withstand the existence of

flaws are measured by tensile testing (ductility), Charpy impact testing (toughness), and COD testing. For welds in FCMs, values of these measurements should be well above the minimum specification requirements. These tests are conducted on the welded joints prepared according to the specified qualification test procedures. An evaluation of ductility, toughness, and crack-opening displacement should be obtained on the weld metal, the HAZ, and the unaffected base metal.


1. R.P. REED, M.B. KASEN, H.I. MCHENRY, C.M. FORTUNKO, AND D.T. REED, FITNESS FOR SERVICE CRITERIA FOR PIPELINE GIRTH WELD QUALITY, BULL. 296, WELDING RESEARCH COUNCIL, JULY 1984 Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Pressure Vessels and Piping

Codes and Specifications The primary users of pressure vessels and piping are the chemical, petroleum, and electric power industries. While there are numerous codes, standards, and specifications that govern the selection of materials for pressure vessels and piping, the manufacture of boilers and pressure vessels in the United States falls under the general specification issued by ASME titled "Boiler and Pressure Vessel Code." This code is published in several parts. The rules governing the design of boilers are included in section I, "Power Boilers." The design rules for pressure vessels are included in section VIII, divisions 1 and 2, "Rules for Construction of Pressure Vessels." The materials requirements are included in section II, "Materials Specifications": • • • •

PART A: FERROUS PART B: NONFERROUS PART C: WELDING RODS, ELECTRODES, AND FILLER METALS PART D: PROPERTIES

The ferrous (SA) and nonferrous (SB) ASME materials specifications were originally developed by ASTM and in most cases are identical. Thus, ASME specifications designated SA 204 and SB 467 may also be found as ASTM specifications A 204 and B 467, respectively. In the United States, the AWS filler metal specifications are the primary guidelines used by the pressure vessel and piping industry. Here again, the ASME specifications (part C) are identical. The ASME specifications carry the initial designation SFA to correspond to the AWS specification designations starting with A; for example, SFA 5.5 is identical to AWS A 5.5. The AWS specifications are written to provide the specific chemical analysis of the filler material and the mechanical properties (minimum tensile and yield strengths) of the deposited weld metal. Standards for pressure piping are covered by the American National Standards Institute (ANSI) document B31.1, "Code for Pressure Piping." This standard also applies to industrial pipelines and piping. Material selection is primarily from the ASTM and ASME materials specifications, although the API does provide a number of pipe specifications for material with yield strengths ranging from 170 to 760 MPa (25 to 110 ksi) and tensile strengths ranging from 275 to 860 MPa (40 to 125 ksi). The intended service temperatures, corrosion resistance, and fabricability are the primary considerations in the selection of materials for pressure vessels and piping. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Environment Ordinary-Temperature Service (-30 to 345 °C, or -20 to 650 °F). The ASME, recognizing the fact that the

ultimate strength of steels remains relatively constant over the temperature range from -30 to 345 °C (-20 to 650 °F), elected to set a safe maximum design stress of 25% of the ultimate tensile strength at temperatures over this temperature range. Consequently, the plain carbon steels are the most economical, with the popular grades being SA 515 ("Pressure

Vessel Plates, Carbon Steel, for Intermediate- and Higher-Temperature Service") and SA 516 ("Pressure Vessel Plates, Carbon Steel, for Moderate- and Lower-Temperature Service"). SA 516, a fine-grain, silicon-aluminum killed steel, is often preferred to SA 515, a coarse-grain, silicon killed steel, because of its better notch toughness characteristics. A steam drum for an electric utility boiler fabricated from SA 515 grade 70 material is shown in Fig. 3. Where reduction in weight is important and operating pressures are high, the higher strength weldable steels are used, such as SA 302 ("Pressure Vessel Plates, Alloy Steel Manganese-Molybdenum and Manganese-Molybdenum-Nickel"). The quenched and tempered version of this steel, SA 533 (grade B, class 1), is used extensively in nuclear reactor vessels. Pertinent data for these materials are listed in Table 5. The specifications listed in Table 5 (and also Tables 6 and 7, to be discussed later) are for plate material. Similar ASME specifications exist for forgings, pipe, tube, and cast material.

TABLE 5 PRESSURE VESSEL STEELS FOR ORDINARY-TEMPERATURE SERVICE ASME NO.

MINIMUM SPECIFIED TENSILE STRENGTH

ALLOY AND NOMINAL COMPOSITION

ALLOWABLE DESIGN STRESS

ksi

mpa

ksi

GRADE

mpa PLAIN-CARBON; COARSE-GRAIN PRACTICE PLAIN-CARBON; FINE-GRAIN PRACTICE; IMPROVED TOUGHNESS 0.26% C MAX; 0.60-0.90% MN 380

55

95

13.75

GRADE

0.27% C MAX; 0.60-0.90% MN

415

60

103

15.00

GRADE

0.29% C MAX; 0.85-1.20% MN

450

65

112

16.25

GRADE 70 SA 302 SA 302 GRADE A GRADE B GRADE C GRADE D SA 202 GRADE A GRADE B SA 537 CLASS 1

0.31% C MAX; 0.85-1.20% MN

485

70

121

17.50

MN-MO

515

75

130

18.8

MN-MO MN-MO-NI MN-MO-NI

550 550 550

80 80 80

138 138 138

20 20 20

CR-MN-SI CR-MN-SI

515 585

75 85

130 147

18.8 21.3

C-MN-SI; NORMALIZED; 63.5 MM (2.5 IN.) 485 MAX THICKNESS C-MN-SI: QUENCHED AND TEMPERED; 63.5 550 MM (2.5 IN.) MAX THICKNESS

70

121

17.5

80

138

20

SA 515 SA 516

55 60 65

CLASS 2

TABLE 6 PRESSURE VESSEL STEELS FOR LOW-TEMPERATURE SERVICE ASME NO.


SA 516

PLAIN-CARBON;

FINE-GRAIN

MINIMUM SPECIFIED TENSILE STRENGTH

ALLOWABLE DESIGN STRESS

LOWEST TEST TEMPERATURE FOR CHARPY V-NOTCH TEST

MPa

MPa

°C

ksi

ksi

°F

GRADE

PRACTICE 0.27% C MAX; 0.60-0.90% MN

415

60

103

15.0

-46

-50

GRADE

0.29% C MAX; 0.85-1.20% MN

450

65

112

16.25

-46

-50

GRADE

0.31% C MAX; 0.85-1.20% MN

485

70

121

17.5

-46

-50

70

...

...

-62

-80

80

...

...

-60

-76

60 65 70 SA 537 CLASS 1

FINE-GRAIN PRACTICE NORMALIZED; 0.24% C MAX; 485 0.70-1.35% MN CLASS QUENCHED AND TEMPERED; 550 0.24% C MAX; 0.70-1.35 MN

2 SA 203 GRADE A GRADE B GRADE D GRADE E SA 645 SA 553 TYPE II TYPE I

2.4% NI

450

65

112

16.25

-60

-75

2.4% NI

485

70

121

17.5

-60

-75

3.5% NI

450

65

112

16.25

-100

-150

3.5% NI

485

70

121

17.5

-100

-150

5.0% NI

655

95

164

23.75

-170

-275

8.0% NI 9.0% NI

690 690

100 100

172 172

25.0 25.0

-170 -195

-175 -320

TABLE 7 PRESSURE VESSEL STEELS FOR HIGH-TEMPERATURE SERVICE ASME NO.


SA 204, GRADE C SA 302, GRADE B SA 387, GRADE 12 CLASS 1

0.28% C MAX; 0.45-0.60% MO

CLASS 2

ALLOWABLE STRESS AT LOW END OF USAGE TEMPERATURE

ALLOWABLE STRESS AT HIGH END OF USAGE TEMPERATURE

°C 430510 0.25% C MAX; 1.15-1.50% 430MN; 0.45-0.60% MO 510

°F 800-950

Mpa 130

Ksi 18.8

Mpa 57

Ksi 8.2

800-950

130

18.8

57

8.2

1.0CR-0.5MO

455565 345480

8501050 650-900

92

13.4

30

43

112

16.3

104

15.1

455565 345480

8501050 650-900

101

14.6

32

4.6

130

18.8

110

15.9

455-

850-

99

14.4

40

5.8

1.0CR-0.5MO

SA 387, GRADE 112 CLASS 1 1.25CR-0.5MO CLASS 2

NORMAL TEMPERATURE RANGE OF USAGE

1.25CR-0.5MO

SA 387, GRADE 22 CLASS 1 2.25CR-1.0MO

CLASS 2

2.25CR-1.0MO

SA 387, GRADE 5 CLASS 1 5.0CR-0.5MO CLASS 2 SA 387, GRADE 9 SA 387, GRADE 91 SA 240 GRADE 304H GRADE 316H GRADE 321H GRADE 347H

5.0CR-0.5MO 9.0CR-1.0MO

595 370480

1100 700-900

480620

510595 9.0CR-1.0MO + NI, V, NB, N, 540AL 650 AUSTENITIC STAINLESS STEELS 18CR-8NI 595815 16CR-12NI-2MO 595815 18CR-10NI-TI 595815 18CR-10NI-NB 595815

119

17.2

109

15.8

9001150

83

12.1

29

4.2

9501100 10001200

73

10.6

22

3.3

99

14.3

28

4.3

11001500 11001500 11001500 11001500

61

8.9

9.7

1.4

71

10.3

8.3

1.2

48

6.9

2

0.3

90

13.0

9

1.3

FIG. 3 STEAM DRUM FOR AN SA 515 STEEL ELECTRIC UTILITY BOILER

Low-Temperature Service (-30 to -195 °C, or -20 to -320 °F). The commonly measured mechanical properties

of steels, such as ultimate strength, yield strength, and elongation, do not always ensure freedom from failure at low temperatures, even when the design load is not exceeded. To ensure safe performance, the steel must be resistant to the initiation and propagation of a crack under all service conditions. Resistance to brittle crack initiation and propagation is commonly referred to as notch toughness or notch ductility. Factors that adversely affect notch toughness include cold working, strain aging, large grain size, increasing section thickness, and increasing the compositional contents of carbon, sulfur, phosphorus, nitrogen, and silicon. Beneficial factors include fine grain size, deoxidation, heat treatment, increasing the manganese and nickel contents, and decreasing the sulfur content. The notch toughness of a steel is also dependent on fabrication procedures, workmanship, and design details. Even the best design will contain some notch effects, and there will also be some imperfections in the highest level of workmanship. In the presence of these notches, the breaking strength of a pressure vessel and, more importantly, its failure mode (ductile or brittle) depend largely on the capacity of the steel at the root of some critical notch to yield plastically as the stress increases. In thick sections, plain carbon steels produced according to fine-grain practice and normalized or quenched and tempered are used for service to -45 °C (-50 °F). Low-carbon high-nickel (e.g., 9% Ni) alloy steels are used for service down to 195 °C (-320 °F). Ferritic steels covered by ASME specifications are shown in Table 6. Austenitic chromium-nickel steels, aluminum, and special copper and aluminum-base alloys have been found to be particularly suitable for applications close to absolute zero. Because austenitic steels have a face-centered cubic (fcc) crystal structure, they retain their toughness to very low temperatures.

High-Temperature Service (345 to 815 °C, or 650 to 1500 °F). In the design of pressure vessels and piping,

engineers and designers are confronted with the problem of selecting materials for a wide range of high-temperature service conditions. The chromium-molybdenum ferritic steels and austenitic stainless steels are generally used for design temperatures above 425 °C (800 °F). In addition to service temperature, corrosion resistance, and fabricability, the following conditions should be considered in high-temperature applications: • • • •

POSSIBLE MAXIMUM TEMPERATURE TYPE AND SIZE OF LOAD EXPECTED LIFE OF THE STRUCTURE COST

Service experience and laboratory test data have established the normal temperature range of usage shown in Table 7 for the materials commonly used for high-temperature service. The allowable stresses in the ASME code, section VIII, division 1, for the low and high ends of the usage temperature ranges are also given in Table 7. The following material properties must be taken into account in addition to the aforementioned service conditions: • • • • • • • • • • •

CREEP STRENGTH STRESS-RUPTURE LIFE DUCTILITY SHORT-TERM TENSILE PROPERTIES SURFACE STABILITY STRUCTURAL STABILITY THERMAL CONDUCTIVITY THERMAL FATIGUE THERMAL EXPANSION WELDABILITY HOT AND COLD WORKABILITY

All of these factors are important in arriving at the most suitable alloy for a given application. Corrosion. Because of the many chemical compounds that may be encountered in pressure vessel and piping service, a

discussion of corrosion is outside the scope of this section. However, there are two areas of concern in both the petrochemical field and the electric power industry: the effect of hydrogen absorption on materials properties and the possible metal loss due to steam/air oxidation. High-Pressure, High-Temperature Hydrogen Service. Pressure vessels and piping exposed to hot high-pressure hydrogen may embrittle, crack, and ultimately fail without visible thinning of the metal or a superficially observable change in appearance of the exposed surfaces. A tool known as the Nelson diagram (Ref 3) is generally used to predict the behavior of alloy material in high-pressure, high-temperature hydrogen (Fig. 4). The Nelson diagram is based on actual service experience and indicates the temperature-pressure limits below which a material has not failed to date. Allowance must be made for temperature and pressure excursions that may occur over the design life of the component. Hydrogen can affect the mechanical properties of steel in two ways. Absorbed hydrogen causes a reduction in yield strength and impairs material toughness. More advanced attack, as manifested by fissuring and decarburization, is reflected in severely reduced tensile strength, ductility, and toughness.

FIG. 4 NELSON DIAGRAM. 1, THE LIMITS DESCRIBED BY THESE CURVES ARE BASED ON ACTUAL SERVICE EXPERIENCE AND ON INFORMATION GATHERED BY API (REF 3). 2, AUTENITIC STAINLESS STEELS ARE GENERALLY NOT DECARBURIZED IN HYDROGEN AT ANY TEMPERATURE OR HYDROGEN PRESSURE. 3, THE LIMITS DESCRIBED BY THESE CURVES ARE BASED ON EXPERIENCE WITH CAST STEEL AS WELL AS ANNEALED AND NORMALIZED STEELS AT STRESS LEVELS DEFINED BY THE ASME CODE, SECTION VIII, DIVISION I. LINES REPRESENT THE UPPER LIMIT OF CONDITIONS FOR ACCEPTABLE USE OF THE ALLOY STEEL. SOURCE:

REF 3

The recent trend in low-alloy steels for hydrogen service has been in the direction of steels of the chromiummolybdenum-vanadium type, such as 3Cr-1Mo-0.25V or 2.25Cr-1Mo-0.25V. For the most severe corrosive conditions, particularly those involving hydrogen sulfide (H2S) in addition to hydrogen, austenitic steels of the 18Cr-8Ni type are used. The National Association of Corrosion Engineers (NACE) standard MR-01-75 is the accepted reference in terms of materials and fabrication requirements for H2S service. This standard specifies the materials and fabrication that are considered to provide satisfactory resistance to H2S attack. Carbon steels for use in wet H2S service are usually evaluated in accordance with NACE TM-01-77 (test solution) and TM-02084 (test method). The relatively high cost of the austenitic steels frequently dictates the use of carbon or low-alloy steel vessels with applied corrosion-resistant cladding. In general, three methods of attaching the cladding to the backing plate are utilized: integral cladding, strip lining, and weld overlay cladding. Integrally clad plate is fabricated by hot rolling in a steel mill or by explosive cladding. Strip lining is applied by using long, narrow alloy sheet attached to the vessel by a continuous weld around the edges. Weld overlay cladding produces a continuous bonded layer of corrosion-resistant alloy on the base metal. Weld overlay cladding can be applied on irregular shapes, such as nozzles, thus avoiding dissimilar-alloy welds in the main pressure boundary. Oxidation in Steam Service. In choosing a steel for service at elevated temperature, consideration must be given not

only to its resistance to applied stress at temperature but also to its resistance to scaling, both of which decrease rapidly with increasing temperature above about 600 °C (1110 °F). Chromium, alone or in combination with silicon or molybdenum, contributes to increased resistance. For plain carbon steels, Fe3O4 is the only oxidation product and is stable until a temperature of 570 °C (1060 °F) is reached. Above 570 °C (1060 °F), FeO is the stable oxide phase and contributes to more rapid metal loss. The addition of chromium raises the temperature at which the Fe3O4/FeO transition occurs (see Table 8). Metal loss due to oxidation is of relatively little importance in heavy-wall pressure vessels; however, it deserves consideration in thinner-wall piping and in thinner low-pressure vessels.

TABLE 8 FE3O4/FEO TRANSITION TEMPERATURES

ALLOY

TEMPERATURE °C °F 0.5CR-0.5MO 590 1095 1CR-0.5MO 600 1110 1.25CR-0.5MO 605 1120 2.25CR-1MO 615 1140 5CR-0.5MO 625 1155 Other Factors Causing Degradation of Materials (Ref 4). Specific service conditions may degrade materials to an

extent that requires replacement or retirement of components in the system. One of the more common failure modes in pressure vessels is SCC, which can be characterized as either intergranular (IGSCC) or transgranular (TGSCC), depending on the environment. Hydrogen-induced cracking (HIC) is another failure mechanism, again influenced by the localized stress condition. Low-alloy steels in service at temperatures above 450 °C (850 °F) will suffer a loss in ductility due to temper embrittlement unless the composition is rigidly controlled. Phosphorus, tin, arsenic, and antimony are the principal elements that contribute to temper embrittlement. The ductility of steels operating in the creep temperature range (generally above 480 °C, or 900 °F) suffers from a condition known as creep embrittlement. The loss in ductility is caused by the formation of small cavities in grain boundaries. In the presence of relatively low stress, the number of cavities grow and eventually link together to form microcracks, thus preventing the plastic deformation that is characteristic of ductile behavior. Sulfur is one of the elements that contribute to cavity formation and loss in ductility in creep service.


3. STEELS FOR HYDROGEN SERVICE AT ELEVATED TEMPERATURES AND PRESSURES IN PETROLEUM REFINERIES AND PETROCHEMICAL PLANTS, 4TH ED., API PUBL. 941, AMERICAN PETROLEUM INSTITUTE, APRIL 1990 4. S. YUKAWA, GUIDELINE FOR PRESSURE VESSEL SAFETY ASSESSMENT, NIST SPEC. PUBL. 780, NATIONAL INSTITUTE FOR STANDARDS AND TECHNOLOGY, P 14-22 Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Fabricability Fabrication operations can affect the performance of pressure vessels and piping in several ways: (1) cold forming operations may lead to strain aging, with a loss in toughness and ductility; (2) hot forming, unless carefully controlled, may produce undesirable changes in microstructure; (3) welding operations may produce an HAZ with reduced toughness and high hardness, and may also produce high tensile residual stresses that could lead to stress-rupture cracking (reheat cracking) during postweld heat treatment of susceptible materials. Welding Process Selection. Considerations which govern the selection of a welding process are those that define the end use of a vessel, such as toughness, tensile properties, creep properties, and corrosion. There are also practical considerations, such as weldment size, equipment limitations, welding position, and access to the weldment area.

Two welding processes are primarily used: shielded metal arc welding (SMAW) and submerged arc welding (SAW). Other processes, such as electroslag welding, gas-tungsten arc welding, and gas-metal arc welding, are also used, but historically SMAW (see the article "Shielded Metal Arc Welding" in this Volume) and SAW (see the article "Submerged Arc Welding" in this Volume) have been dominant in vessel manufacture. A characteristic of both the SMAW and SAW processes is that they rely on the use of a flux to produce a slag over the molten weld pool, protecting it from atmospheric contamination. These fluxes vary significantly in formulation and operating characteristics and can modify the chemical composition, inclusion content, and microstructure of the final weld metal. Weldability. Steels for pressure vessel fabrications are often classified as weldable based on composition, thickness, and

need for preheat (see the article "Weldability Testing" in this Volume). Carbon and low-alloy steels are the most frequently used materials for pressure vessels; consequently, their weldability has received the greatest attention. Preheating is utilized to prevent cracking, to permit diffusion of hydrogen from the weld and HAZ during and following welding, to reduce distortion, and to prevent loss of toughness and ductility. Preheating is also used to ensure that all traces of moisture have been removed from the surface of the material. Of the various schemes proposed for predicting preheating requirements, most are based on the carbon equivalent (Ceq). A variety of equations have been developed by various researchers for determining the Ceq, one of which has demonstrated that prevention of cracking due to hydrogen can be correlated with chemical composition in terms of a carbon equivalent formula of the form: Ceq = C +

Mn Si Ni Cr Mo + + + + 6 24 40 5 4

This formula was developed by Winn (Ref 5). A plot of carbon equivalents for a series of steel compositions used by Winn against preheat temperature is shown in Fig. 5. A linear relationship exists between the steel chemical composition, expressed by the carbon equivalent formula, and the minimum preheat temperature to avoid HAZ cracking.

FIG. 5 CARBON EQUIVALENT VERSUS MINIMUM PREHEAT TEMPERATURE. THE BEST-FIT LINE SHOWN MAY BE REPRESENTED APPROXIMATELY BY: T = 210(CEQ) - 25, WHERE T IS THE PREHEAT TEMPERATURE (°C).

The formula used by the International Institute of Welding (IIW) for judging the risk of underbead cracking and the resultant need for preheat is as follows: Ceq = C +

Mn Cr + Mo + V Ni + Cu + + 6 5 15

Regardless of the formula used, the Ceq values must be considered as approximations, because all factors affecting hardenability (e.g., section thickness, mass, restraint, grain size, and environmental effects) are not included. Weld Zones. Based on work published by Savage, Nippes, and Szekeres (Ref 6), a typical weld consists of the

composite zone, the unmixed zone, the weld interface, the partially melted zone, the HAZ, and the unaffected base metal (Fig. 6). All of these zones possess different mechanical properties, which are determined by the chemical compositions of the weld metal and the base metal. The preheat used, the weld energy input, and whether or not postweld heat treatment is specified also affect the resultant mechanical properties of the weldment.

FIG. 6 TYPICAL ZONES ASSOCIATED WITH FUSION WELDS. SOURCE: REF 6

Postweld Heat Treatment. The circumstances under which the postweld heat treatment of pressure vessels is required

are set forth in the ASME pressure vessel codes. They are determined primarily by the steel composition, thickness, specific fabrication operations, and intended service. The principal beneficial effects of postweld heat treatment are the relief of residual stresses and the tempering action on the microstructures produced by the welding thermal cycles. Undesirable metallurgical phenomena may also be produced, such as temper embrittlement or cracking due to stress-rupture failures at grain boundaries.


5. W.H. WINN, WELDABILITY OF LOW ALLOY STEELS, BR. WELD. J., VOL 11 (NO. 8), AUG 1964 6. W.F. SAVAGE, E.F. NIPPES, AND E.S. SZEKERES, A STUDY OF WELD INTERFACE PHENOMENA IN A LOW ALLOY STEEL, WELD. J., SEPT 1976, P 260S-268S Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Mechanical Properties of Weldments Design Considerations. Some designers have tacitly assumed that a pressure vessel behaves as if it were manufactured from an isotropic material with the same properties throughout the vessel. Because of the presence of welded seams and welded attachments, however, no actual vessel fulfills this simple concept. Welds produce regions where the mechanical properties are considerably different from those of the base material. The stress-rupture life and the ductility at rupture may also be significantly different. Weld Metal Strength. There is a tendency for weld metal to overmatch the parent metal in strength at ambient and low

temperatures. While a slight overmatch is not harmful, excessive hardness in welds can cause greater susceptibility to weld cracking. Creep of Welds in High-Temperature Service. Recent developments suggest that welded joints in some chromium-

molybdenum steels are a potential source of weakness in creep. Vessels that have performed satisfactorily in service thus far may have safety factors lower than those suggested by present design rules. This arises from the mismatch of creep

properties in the different weld zones, especially in thick sections. In recognition of this fact, section III of the ASME code, "Nuclear Power Plant Components," has assigned stress-reduction factors to weldments operating in the creep range of temperatures. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Shipbuilding and Offshore Structures

The materials used for ships and offshore structures encounter a broad range of service conditions and must be capable of withstanding the rigors of weather and sea. With the advent of sophisticated communications systems, ships can often avoid heavy weather, whereas site-dependent offshore structures must be built to withstand any service condition that arises. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Ships Steel ship structures are comprised of shell plating and shapes joined by welding. The structures vary considerably among various ship types, including small boats 10 to 25 m (30 to 80 ft) long to supertankers with lengths ranging from 100 to 400 m (330 to 1300 ft) or more. In between are container ships, liquefied natural gas (LNG) and liquefied petroleum gas (LPG) carriers, passenger ships, ore carriers, and others. Requirements for ordinary- and higher-strength steels according to the rules of the American Bureau of Shipping Rules (ABS) are given in Tables 9 and 10, respectively.

TABLE 9 ABS REQUIREMENTS FOR ORDINARY-STRENGTH HULL STRUCTURAL STEEL ( ≤ 50 MM, OR 2 IN., THICK)

CHEMICAL COMPOSITION(A), % C Mn P S Si A SEMIKILLED OR KILLED 0.23 2.5 × C MIN FOR 0.04 0.04 ... (B) FOR MATERIAL THICKER MAX PLATE THICKER MAX MAX THAN 12.5 MM (0.5 IN.) THAN 12.5 MM (0.5 IN.) B SEMIKILLED OR KILLED 0.21 0.08-1.10; 0.60 MIN 0.04 0.04 0.35 MAX FOR FULLY KILLED MAX MAX MAX OR COLD FLANGING D FULLY KILLED; FINE0.21 0.70-1.35; 0.60 MIN 0.04 0.04 0.10GRAIN PRACTICE(C) MAX FOR THICKNESS ≤ 25 MAX MAX 0.35 MM (1 IN.) E FULLY KILLED; FINE0.18 0.70-1.35 0.04 0.04 0.10GRAIN PRACTICE MAX MAX MAX 0.35 TENSILE STRENGTH YIELD POINT (MIN) ELONGATION (MIN) Mpa Ksi Mpa Ksi FOR ALL GRADES: FOR ALL GRADES: FOR ALL GRADES: 400-490 58-71 235 34 21% IN 200 MM (8 IN.) OR 24% IN 50 MM (2 IN.) FOR GRADE A SHAPES FOR GRADE A THICKER THAN FOR COLD FLANGING QUALITY: GRADE DEOXIDATION

AND BARS: 400-550 58-80 FOR COLD FLANGING QUALITY: 380-450 55-65 GRADE

25 MM (1 IN.) 220 32 FOR COLD FLANGING: 205

23% MIN IN 200 MM (8 IN.)

30

CHARPY VNOTCH IMPACT TEMPERATURE °C °F

ENERGY AVERAGE (MIN), J (FT · LBF)

LONGITUDINAL SPECIMENS

TRANSVERSE SPECIMENS

A B D

... 0 -10

... 32 14

... 27 (20) 27 (20)(D)

... 20 (14) 20 (14)(D)

E

-40

-40

27 (20)

20 (14)

HEAT TREATMENT

MARKING

... ... NORMALIZED FOR MATERIAL THICKER THAN 35 MM (1.4 IN.)(E) NORMALIZED(G)

AB/A AB/B AB/D(F)

AB/E

(A) FOR ALL GRADES EXCLUSIVE OF GRADE A SHAPES AND BARS, THE SUM OF CARBON CONTENT AND ONE-SIXTH OF THE MANGANESE CONTENT IS NOT TO EXCEED 0.40%. THE UPPER LIMIT OF MANGANESE MAY BE EXCEEDED UP TO A MAXIMUM OF 1.65%, PROVIDED THAT THIS CONDITION IS SATISFIED. THE CONTENTS OF NICKEL, CHROMIUM, MOLYBDENUM, AND COPPER ARE TO BE DETERMINED AND REPORTED, AND WHEN THE AMOUNT PRESENT IS LESS THAN 0.02%, THESE ELEMENTS MAY BE REPORTED AS 0.02%. (B) A MAXIMUM CARBON CONTENT OF 0.26% IS ACCEPTABLE FOR GRADE A PLATES 12.5 MM (0.5 IN.) AND ALL THICKNESSES OF GRADE A SHAPES AND BARS. (C) GRADE D MAY BE FURNISHED SEMIKILLED IN THICKNESSES UP TO 35 MM (1.4 IN.), PROVIDED THAT STEEL >25 MM (1 IN.) THICK IS NORMALIZED. IN THIS CASE THE REQUIREMENTS RELATIVE TO MINIMUM SILICON AND ALUMINUM CONTENTS AND FINEGRAIN PRACTICE DO NOT APPLY. (D) IMPACT TESTS ARE NOT REQUIRED FOR NORMALIZED GRADE D STEEL WHEN FURNISHED FULLY KILLED, FINE-GRAIN PRACTICE. (E) CONTROLLED ROLLING OR THERMOMECHANICAL CONTROLLED ROLLING OF GRADE D STEEL MAY BE CONSIDERED AS A SUBSTITUTE FOR NORMALIZING. (F) GRADE D HULL STEEL THAT IS NORMALIZED, THERMOMECHANICAL CONTROLLED ROLLED, OR CONTROLLED ROLLED IN ACCORDANCE WITH (E) IS TO BE MARKED AB/DN. (G) CONTROLLED ROLLING OR THERMOMECHANICAL CONTROLLED ROLLING OF GRADE E SHAPES AND THERMOMECHANICAL CONTROLLED ROLLING OF GRADE E PLATES MAY BE CONSIDERED AS A SUBSTITUTE FOR NORMALIZING.

TABLE 10 ABS REQUIREMENTS FOR HIGHER-STRENGTH HULL STRUCTURAL STEEL ( ≤ 50 MM, OR 2 IN., THICK) GRAD E(A)

AH3 2

DEOXIDA TION

SEMIKI LLED

TENSIL E STREN GTH

YIELD POINT (MIN)

Mp a

Ks i

M pa

47 0-

6 31 45 8- 5 .5

Ksi

ELONGA TION (MIN)

CHARPY VNOTCH IMPACT TEMPERA TURE °C °F

ENERGY AVERAGE (MIN), J (FT · LBF)

LONGITU DINAL SPECIMEN S

TRANSV ERSE SPECIM ENS

0

34 (25)(B)

24 (17)(B)

32

HEAT TREATMENT

MARKIN G

...

AB/AH 32

OR KILLED AH3 6

DH3 2

DH3 6

EH3 2

EH3 6

KILLED; FINEGRAIN PRACTI CE(C)

KILLED; FINEGRAIN PRACTI CE (C)

58 5 49 062 0 47 058 5 49 062 0

8 5 7 19 0 6 88 5 7 19 0

47 058 5

6 31 45 8- 5 .5 8 5

49 062 0

7 35 51 1- 5 9 0

35 51 5

31 45 5 .5

35 51 5

AB/AH 36

-20 FOR ALL GRADE S: 19% IN 200 MM (8 IN.) OR 22% IN 50 MM (2 IN.) -40

-4

-40

34 25(B)

34 (25)

24 (17)(B)

24 (17)

NORMALI ZED(D) FOR MATERIA L >12.5 MM (0.5 IN.) THICK(E)

AB/DH 32(F)

NORMALI ZED(G)

AB/EH 32

AB/DH 36(F)

AB/EH 36

(A) THE NUMBERS FOLLOWING THE GRADE DESIGNATION INDICATE THE YIELD POINT OR YIELD STRENGTH TO WHICH THE STEEL IS ORDERED AND PRODUCED IN KGF/MM2. THE CHEMICAL COMPOSITION FOR ALL GRADES IS: 0.18% MAX C; 0.90-1.60% MN (GRADE AH THAT IS ≤ 12.5 MM, OR 0.5 IN., THICK MAY HAVE A MINIMUM MANGANESE CONTENT OF 0.70%); 0.04% MAX P; 0.04.% MAX S; 0.10-0.50% SI (GRADE AH PLATE THAT IS ≤ 12.5 MM, OR 0.5 IN., THICK AND ALL THICKNESSES OF AH SHAPES MAY BE SEMIKILLED, IN WHICH CASE THE 0.10% MIN SI DOES NOT APPLY; UNLESS OTHERWISE SPECIALLY APPROVED, AH PLATE >12.5 MM, OR 0.5 IN., THICK ARE TO BE KILLED WITH 0.10-0.50% SI); 0.40% MAX NI; 0.25% MAX CR; 0.08% MAX MO: 0.35% MAX CU; 0.05% MAX NB; 0.10 MAX V. NICKEL, CHROMIUM, AND MOLYBDENUM MAY BE REPORTED AS 0.02% WHEN THE AMOUNT PRESENT IS LESS THAN 0.02%. NIOBIUM AND VANADIUM NEED NOT BE REPORTED ON THE MILL SHEET UNLESS INTENTIONALLY ADDED. (B) IMPACT TESTS ARE NOT REQUIRED FOR AH THAT IS ≤ 12.5 MM (0.5 IN.) THICK OR FOR ALUMINUM TREATED GRADE AH THAT IS ≤ 35 MM (1.4 IN.) THICK, NOR FOR FULLY KILLED FINE-GRAIN NORMALIZED GRADES AH OR DH THAT IS ≤ 50 MM (2 IN.) THICK. (C) GRADES DH AND EH ARE TO CONTAIN AT LEAST ONE OF THE GRAIN-REFINING ELEMENTS IN SUFFICIENT AMOUNT TO MEET THE FINE-GRAIN PRACTICE REQUIREMENT. (D) CONTROLLED ROLLING OR THERMOMECHANICAL CONTROLLED ROLLING OF GRADE D STEEL MAY BE CONSIDERED AS A SUBSTITUTE FOR NORMALIZING. (E) CONTROLLED ROLLING OR THERMOMECHANICAL CONTROLLED ROLLING OF GRADE EH SHAPES AND THERMOMECHANICAL CONTROLLED ROLLING OF GRADE EH PLATES MAY BE CONSIDERED AS A SUBSTITUTE FOR NORMALIZING. (F) THE MARKING AB/DHN IS TO BE USED TO DENOTE GRADE DH PLATES WHICH HAVE EITHER BEEN NORMALIZED, THERMOMECHANICAL CONTROLLED ROLLED, OR CONTROL ROLLED IN ACCORDANCE WITH APPROVED PROCEDURE. (G) FOR DH TREATED WITH NIOBIUM OR VANADIUM, OR WHEN USED IN COMBINATION WITH ALUMINUM OR EACH OTHER. DH TREATED ONLY WITH ALUMINUM IS TO BE

NORMALIZED WHEN >25 MM (1 IN.) THICK. Tankers and Container Ships. Large ocean-going tankers have been built with both single and double hulls, with the latter type becoming more popular due to environmental considerations and the belief that oil spills are more readily prevented by double-hull construction. However, most tankers in service today have single hulls. In the design of typical ships, 0 °C (32 °F) is normally assumed to be the lowest service temperature.

Material selection for single-hull tankers for a particular steel strength level is based on stress level, location, thickness, and toughness. A typical cross section of the shell and longitudinal bulkheads for a large tanker is shown in Fig. 7. Plating for various parts of the ship can be categorized as secondary, intermediary, and special application. The four corners of a ship and the deck and bottom plating in the way of the longitudinal bulkheads are usually highly stressed and are made from "special" steels with improved toughness properties. The side shell and areas close to the neutral axis of the ship are the least stressed, and thus the toughness of the steels employed is of "secondary" importance. The remaining deck and bottom of a tanker are moderately stressed and are constructed of intermediate-toughness steels.

FIG. 7 MIDSHIP CROSS SECTION OF A LARGE TANKER, INDICATING SHELL PLATING AND LONGITUDINAL BULKHEADS.

The design considerations for tankers can also apply to container ships. However, because of the large cutouts on the deck that allow access for loading and unloading, the most critical structural area of a container ship is the box girder that connects the side shell to the deck. Major stresses are transmitted through this member, and thus the steels selected for these areas must have improved notch toughness. Liquefied Natural Gas/Liquefied Petroleum Gas Ships. The minimum service temperature for most ships is

approximately 0 °C (32 °F), and hull steels are selected and applied on this basis. Where temperatures below 0 °C (32 °F) are expected, application adjustments or selection of materials with superior low-temperature notch toughness must be considered. In the marine industry, low-temperature service is a significant factor in the design of ships carrying liquefied gases in bulk and of structures operating for prolonged periods in cold weather. Liquefied gases--both petroleum and natural--are transported in bulk by ship at temperatures ranging from -50 to -165 °C (-58 to -265 °F). Steels for ship hulls are brittle at cryogenic temperatures. Thus, for safety reasons, it is essential to prevent the liquefied gas from contacting the steel hull and to minimize excessive conduction cooling of the hull. The primary requisite for the cargo tank material is high notch toughness at low temperature. Also, because shipboard tanks are subjected to dynamic loads resulting from ship motions, fatigue is another consideration.

Fracture mechanics concepts, which require knowledge of fracture toughness and fatigue crack-growth rate properties at low temperature, are frequently used in design. One design concept applied is "leak before failure," combined with a small-leak protection system. Utilizing this concept, a ship will be able to journey to an unloading and repair site before a detected flaw grows to critical size. Other systems are provided with large secondary containment barriers that can contain major leaks while permitting the ship to proceed to an unloading facility. All designs incorporate insulation to prevent excessive conduction cooling of the hull structure. For LNG containment systems, designed for cargo temperatures of -65 °C (-265 °F), external insulations of balsa, pearlite, polyurethane foam, and polyvinyl chloride foam have been used. For higher-temperature liquefied gases, such as LPG transported at a temperature of about -50 °C (-60 °F), some tanks have been internally insulated, which permits the use of less expensive tank structure materials than would be necessary with external insulation. The worldwide adoption of International Maritime Organization (IMO) requirements for the carriage of liquefied gases by sea has unified material requirements for ships carrying cryogenic cargos in bulk. For the primary containment of such cryogenic cargos as LNG (methane) at a temperature of -165 °C (-265 °F), the most commonly selected materials have hexagonal close-packed (hcp) crystal structures; these include alloys such as 5083 aluminum, austenitic stainless steels, and 36% Ni steels, selected because of their superior low-temperature toughness properties. Nickel-bearing body-centered cubic (bcc) alloys, such as 9% Ni steels, have also been used, although less extensively. The most important characteristic of the low-temperature mechanical behavior of ferritic steels is the ductile-to-brittle fracture transition. The transition occurs over a narrow temperature range for sharply notched specimens, and at temperatures below the transition, failure occurs by the brittle cleavage mode. This low-temperature embrittlement severely restricts the use of ferritic materials in cryogenic systems. Structural applications of ferritic materials are generally limited to environments where the service temperature remains above the ductile-to-brittle transition. The ferritic steels that are used for low-temperature applications usually contain nickel to improve low-temperature toughness. At the lower end of the cargo temperature range, austenitic alloys and 5083 aluminum are commonly used. Table 11 indicates the service temperature specified in the ABS rules for each material. Austenitic stainless steel and 36% Ni steel are used for nonpressurized membrane tank designs. These alloys are provided in thin 0.5 to 1.5 mm (0.02 to 0.06 in.) sheets, usually corrugated to provide stiffness and allow for thermal expansion.

TABLE 11 MATERIALS FOR CRYOGENIC APPLICATIONS

SERVICE TEMPERATURE °C °F 2% NI STEEL A 203, GRADE A -62 -80 A 203, GRADE B -59 -75 3% NI STEEL A 203, GRADE D -90 -130 A 203, GRADE E -79 -1L0 (A) 5% NI STEEL A 645 -105 -155(A) 9% NI STEEL A 353 -196 -320 9% NI STEEL A 553 -196 -320 AUSTENITIC STAINLESS STEEL A 240 -196 -320 (B) 36% NI ALLOY A 658 -196 320(B) ALUMINUM ALLOY E 209, ALLOY 5083 -196 -320 MATERIAL

ASTM NO.

(A) 5% NI STEEL MAY BE USED DOWN TO -165 °C (-265 °F) PROVIDED THAT IMPACT TESTS ARE CONDUCTED AT -196 °C (-320 °F). (B) CHEMISTRY WILL BE SPECIALLY CONSIDERED FOR LOWERING THE COEFFICIENT OF EXPANSION. For spherical pressurized tanks, 5083 aluminum alloy at a thickness of 40 to 60 mm (1.6 to 2.4 in.) is used to form tanks up to 35 m (115 ft) in diameter. The aluminum tank is attached to the steel hull by means of an explosion-bonded aluminum-to-steel plate; the tank connects to the aluminum side, and the steel side connects to the hull. This type of explosion-bonded plate is also used for connecting aluminum superstructures to steel decks.


Offshore Structures Several important characteristics and operations of many offshore drilling units dictate the use of materials different from those used for ships, or restrictions on the application of similar materials. Much offshore exploration by mobile drilling or fixed units is carried out in areas where ambient air temperatures remain below 0 °C (32 °F) for long periods, such as in the Arctic. Toughness for these applications must be considered on the basis of the lowest expected steel temperature, because ferritic steels undergo a ductile-to-brittle fracture transition. The lowest steel temperature expected in service is estimated from meteorological data representing the coldest region in which the structure is designed to operate. Toughness Versus Structural Application. In general, toughness testing of structural steel for offshore structures is conducted at a temperature below the minimum expected service temperature, the extent of the temperature increment being based on certain design criteria, such as anticipated stresses, stress concentrations, and redundancy of the structural member. For example, the ABS defines three structural categories to which different levels of toughness are assigned (see Table 12). The three categories are assigned to structural components based on consideration of the following: consequence of failure; redundancy of component; service experience with similar design; stress level, including stress concentration and restraint; fatigue; loading rate; and capability of inspection and repair in service.

TABLE 12 TEST TEMPERATURE FOR STEELS WITH A YIELD STRENGTH OF LESS THAN 415 MPA (60 KSI)

STRUCTURAL DESCRIPTION CATEGORY SECONDARY LEAST CRITICAL: REDUNDANT PRIMARY

SPECIAL

MAIN STRUCTURAL MEMBERS THAT SUSTAIN PRIMARY TENSILE STRESS AND WHOSE FAILURE WOULD JEOPARDIZE THE SAFETY OF THE STRUCTURE MOST CRITICAL: FRACTURE CRITICAL MEMBERS THAT COULD EXPERIENCE RAPID LOADING AT POINTS OF STRESS CONCENTRATION

CHARPY V-NOTCH TEST TEMPERATURE AT SERVICE TEMPERATURE 10 °C (18 °F) BELOW SERVICE TEMPERATURE 30 °C (55 °F) BELOW SERVICE TEMPERATURE

Absorbed energy obtained by a Charpy V-notch test is the most commonly used criterion for grading steels according to toughness. This criterion cannot be related directly to design, but is used as an expression of a toughness quality level. Structural steel elements for offshore structures are often fabricated from cold-formed sections. As most structural steels are strain-age sensitive, it is important to evaluate the extent to which a given steel loses toughness after being cold strained. Outer fiber strains of less than 3% usually do not reduce toughness significantly. In general, at above 3% strain the material is either heat treated to restore the toughness or tested to determine that the poststrain toughness has not degraded below minimum acceptable values. Quenched and Tempered Steels. Certain offshore drilling unit designs, especially the self-elevating type, make

weight savings in the upper structure and the legs an important consideration in order to reduce the size and power requirements for lifting equipment and to aid stability. Therefore, high-strength quenched and tempered martensitic steels with yield strengths of approximately 700 MPa (100 ksi) find wider application in offshore structures than in ships. Areas subject to concentrated loads, such as openings for leg structures, cantilever overhangs for drilling booms, helicopter decks, and crane pedestals, are other locations on offshore drilling structures in which high-strength steels are commonly used.


Welding Process Selection Welding has exclusively replaced riveting for joining structural steels in the fabrication of ships and offshore structures. In shipbuilding, welding permits large subassemblies to be built quickly and efficiently; these subassemblies are then welded together to form ship hulls. Manual Welding. Shielded metal arc welding with stick electrodes, semiautomatic gas metal arc welding (GMAW), and flux-cored arc welding (FCAW) are all used in shipbuilding. These processes offer a high degree of versatility because of the portability of equipment and are commonly used for short weld runs and for erection welds. Gravity-feed welding is used in the assembly of ships and offshore structures. It is carried out using contact or drag-

type electrodes with tripod-type holding mechanisms, which allow welding to be done semiautomatically. An operator places the electrode in the holder and positions the holding mechanism, and welding progresses unrestricted until the electrode is consumed. Long electrodes enable welds more than 750 mm (30 in.) in length to be made with one electrode. Specific electrode diameters with proper current and voltage settings control the size of the deposited weld. One operator can handle two to six gravity-type devices at the same time. This process has been largely employed for attaching stiffeners to plating by fillet welding. It is particularly attractive when the weld lengths are less than 3 m (10 ft), in which case the use of a fully automatic process may not be justified due to setup time, portability, and other considerations. Applications include internal stiffening of drilling-rig support columns and braces, as well as stiffening of decks. Mechanized Welding. Submerged arc welding is used for making long, continuous butt welds to join plate panels or for

joining stiffeners to plating by fillet welding. Modern shipyards that build large ships, especially tankers, where a considerable amount of flat panel work is required, have invested heavily in panel lines to further mechanize or automate the SAW process. Such panel lines are capable of making panel subassemblies up to about 20 m2 (200 ft2) in size. One-side welding using SAW was introduced and perfected in Japan during the 1960s and early 1970s in connection

with a massive tanker-building program for very large and ultralarge crude-oil carriers. In one-side welding, weld backing on the underside is provided by copper bars, ceramic forms, solid flux systems, or powdered flux held in place by various troughing systems. Electroslag welding (ESW) and electrogas welding (EGW) have been used to make relatively long (more than 10

m, or 30 ft) continuous vertical erection welds in the side shells of large ships. These processes offer higher deposition and production rates than more conventional welding processes. The minimum plating thickness is about 19 to 25 mm (

3 4

to 1 in.). Consumable-nozzle ESW, a variation of the basic electroslag process, is used for making vertical butt welds in deck and bottom longitudinals and in heavy face plates of girders. Oxyfuel Gas (OFG) and Plasma Arc Cutting. Oxyfuel gas cutting is used predominantly for cutting panels and

stiffeners and for handfitting miscellaneous members. Plasma arc cutting has been used for mechanized cutting of plating up to about 12 to 20 mm (0.5 to 0.8 in.) thick. It offers the advantage of faster cutting speeds and less heat distortion as compared with gas cutting for thin plating. Because the significant tonnages of plating and shaping that must be cut or beveled for welding large ships, a high degree of mechanization has been introduced. Numerical control cutting permits panels up to about 2.5 by 15 m (8 by 50 ft) to be completely cut to size with appropriate cutouts, as well as edge beveling for welding to make large subassemblies. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Weld Considerations The materials used for ship construction are generally ordinary-strength or higher-strength steels with a maximum yield strength of 355 MPa (52 ksi) (see Tables 9 and 10). For offshore drilling units and platforms, a variety of high-strength quenched and tempered steels with yield strengths up to 700 MPa (100 ksi) are also used. Weld filler metals should approximate the tensile and toughness properties of the base material; where materials of two different strengths or toughnesses are to be joined, the lower-strength or lower-toughness weld metal can be used. Weld metal toughness is an important parameter in shipbuilding. Quantitative measurements of toughness are based on Charpy V-notch properties, and values comparable to those of the base materials are specified for weld metal. For offshore structures such as fixed platforms or drilling units, which must stay on station and encounter severe storms, winds, waves, and temperature fluctuations, toughness is also an important consideration. Materials and weld metals with yield strengths of up to 700 MPa (100 ksi) are commonly offered for tubular leg chord sections, braces, and rack material. The racks that form part of the leg chord are often 100 to 150 mm (4 to 6 in.) thick. Hydrogen Cracking. When joining thick high-strength quenched and tempered steels, appropriate preheat and interpass

temperatures are necessary to avoid HIC (see the article "Cracking Phenomena Associated with Welding" in this Volume). This condition is generally more prevalent in HSLA steels than in ordinary-strength steels, and its occurrence is generally proportional to the hardenability and tensile strength of the base material. Hydrogen cracks are usually found in the HAZ and can occur in both fillet- and butt-welded connections. Hydrogen cracking is often referred to as cold or delayed cracking because of the tendency for these cracks to occur and propagate after the weld has cooled to ambient temperatures. It is not uncommon for such cracks to be revealed hours or even days after welding, depending on the restraint conditions within the structure. For this reason, it is common practice to postpone the final nondestructive testing until several days after completing a weld. Hydrogen sources in and around the arc atmosphere include moisture in the electrode coating, flux, gas, and base material and the presence of hydrocarbons, such as grease, paint, or oil. Moisture requirements for electrode coatings, along with storage and conditioning recommendations, are included in specifications and become more severe as strength levels increase. In SAW, where the flux is a common source of moisture, storage should be controlled to avoid moisture contamination from the atmosphere. Inadequate protection from hydrogen must be provided, especially in final erection or field welding applications, where hydrogen sources are more prevalent and where accessibility to the welded joint is less than ideal. Lamellar tearing is the separation of steel plate underneath a weld in a plane generally parallel to the plate surface. It

usually results from high through-thickness strains induced by weld shrinkage. These strains cause nonmetallic inclusions present in the steel to decohere. The resulting tears or cracks are planar in appearance, with some steplike characteristics. Welded connections joining two or more tubular sections are common in offshore structures. On several occasions, lamellar tearing has been observed in diaphragm plates used between two tubulars to simplify construction. Improvement in through-thickness properties is possible by using steels with fewer impurities or nonmetallic inclusions, lowering the permitted sulfur contents (typically 0.010% maximum), and treating the molten steel with rare earth elements, which combine with sulfur to form refractory sulfides that resist deformation at hot rolling temperatures. Another approach to eliminate lamellar tearing, and to simplify fabrication, is to use steel castings for the node connections. Centrifugally cast straight tubulars have also been proposed for offshore structural applications. Lamellar tearing has been reported in sheer-strake-to-deck connections in ships. This problem can be eliminated by the use of gunwale connections. Lamellar tearing or cracking at welded connections is best avoided by the selection of a design that reduces the tendency for stressing through the thickness of the structural members. In the construction of offshore drilling units and platforms, lamellar tearing presents a more acute problem because of the complexity of welded joints and connections. Sometimes joints may be unavoidably susceptible to lamellar tearing; in such cases, carefully planned welding sequences should be employed. It is common practice to use other than fullpenetration welds; for instance, double-bevel partial-penetration T-welds are often used. Steel exhibiting the equivalent of 20% (minimum) reduction of area in a through-thickness tensile specimen (as per ASTM A 770) is considered to be less susceptible to lamellar tearing. Commonly used structural steels, including ABS hull steels, that meet the A 770 specification are available from some producers. Such steels have additional restrictions on chemistry and are produced using special melting practices.

Undermatching Weld Metal Strength. In some cases, weld metal that undermatches the base plates may be used in

the fabrication and repair of high-strength steels, primarily for fillet welding and local repairs. The use of the lowerstrength weld metal minimizes cracking tendency by providing a weld with reduced residual stresses and somewhat higher ductility. Such lower-strength filler metals are permitted when the strength level of the particular weld fulfills design requirements. When lower-strength covered electrodes are used to weld higher-strength materials, special precautions associated with welding the higher-strength materials are still required to prevent hydrogen cracking. When E80xx or E90xx electrodes are used to weld 700 MPa (100 ksi) yield strength material, such as ASTM A 514, the electrodes should be subjected to appropriate baking to reduce electrode-covering moisture to 0.15% maximum. This precaution should also be taken in the more common case where E7018-type electrodes are used to join high-strength steels (e.g., 700 MPa, or 100 ksi, yield strength) to lower-strength steels. To summarize, in the following circumstances, the use of undermatching electrodes should be considered: •

•

FILLET WELDING: USE OF E9018 OR EVEN E8018 ELECTRODES FOR FILLET WELDING OF STEELS WITH A YIELD STRENGTH OF 700 MPA (100 KSI) IS A CONVENIENT WAY OF MINIMIZING CRACKING TENDENCY IN THE WELD METAL AND HAZ. ELECTRODES SHOULD BE BAKED TO A 0.15% MAXIMUM MOISTURE CONTENT. REPAIR WELDING: THE LOWER STRENGTH AND LOWER RESIDUAL STRESS ASSOCIATED WITH THE E9018 ELECTRODE MAKE IT THE ELECTRODE OF CHOICE FOR ISOLATED REPAIRS IN BUTT OR FILLET WELDS IN THE HIGHER-STRENGTH STEELS (700 MPA, OR 100 KSI, YIELD STRENGTH), PROVIDED THAT THE OVERALL STRENGTH OF THE STRUCTURE IS MAINTAINED TO MEET DESIGN REQUIREMENTS.


Aerospace

Aerospace systems push materials to the limits of their physical and mechanical properties. Therefore, it is important to optimize designs, which often involves consideration of nonmechanical fastening methods. Welding, brazing, soldering, and/or adhesive bonding are the only options in many aerospace structures. Whether the driving force is weight, leakproof seals, continuous load paths, or any number of other requirements, nonmechanical joining of materials is essential in the manufacture of aerospace systems. The list of materials used in the construction of such systems is quite large. This section will present an overview of mechanical and physical property considerations as they relate to joining. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Material Selection Criteria Properties. Modern aerospace systems employ metals, ceramics, composites, and polymers. Important mechanical and

physical properties for aerospace applications are listed in Table 13 (Ref 7).

TABLE 13 MATERIAL PROPERTIES OF IMPORTANCE IN AEROSPACE APPLICATIONS

PHYSICAL PROPERTIES

CRYSTAL STRUCTURE DENSITY MELTING POINT VAPOR PRESSURE VISCOSITY POROSITY PERMEABILITY REFLECTIVITY TRANSPARENCY OTHER OPTICAL PROPERTIES DIMENSIONAL STABILITY ELECTRICAL/MAGNETIC PROPERTIES CONDUCTIVITY DIELECTRIC CONSTANT COERCIVE FORCE HYSTERESIS SUSCEPTIBILITY PERMEABILITY REMANENCE MECHANICAL PROPERTIES HARDNESS MODULUS OF ELASTICITY TENSION COMPRESSION POISSON'S RATIO STRESS-STRAIN CURVE YIELD STRENGTH TENSION COMPRESSION SHEAR ULTIMATE STRENGTH TENSION SHEAR BEARING FATIGUE PROPERTIES SMOOTH NOTCHED CORROSION FATIGUE ROLLING CONTACT FRETTING CHARPY TRANSITION TEMPERATURE FRACTURE TOUGHNESS (KIC) HIGH-TEMPERATURE PROPERTIES CREEP STRESS RUPTURE DAMPING PROPERTIES WEAR PROPERTIES GALLING ABRASION EROSION CAVITATION

SPALLING BALLISTIC IMPACT THERMAL PROPERTIES CONDUCTIVITY SPECIFIC HEAT COEFFICIENT OF EXPANSION EMISSIVITY ABSORPTIVITY ABLATION RATE FIRE RESISTANCE MAXIMUM/MINIMUM OPERATING TEMPERATURE CHEMICAL PROPERTIES POSITION IN ELECTROMOTIVE SERIES CORROSION AND DEGRADATION ATMOSPHERIC SALT WATER ACIDS HOT GASES ULTRAVIOLET OXIDATION THERMAL STABILITY BIOLOGICAL STABILITY STRESS CORROSION HYDROGEN EMBRITTLEMENT HYDRAULIC PERMEABILITY FABRICATION PROPERTIES CASTABILITY HEAT TREATABILITY HARDENABILITY FORMABILITY MACHINABILITY WELDABILITY Source: Ref 7 Applications. The service requirements for a particular application will order the material properties. As shown in Fig. 8,

material selection is based on a number of interrelated factors, including material properties, manufacturability, repairability, and availability (Ref 8). Because compromise is required to ensure optimal material selection, both the service requirements and the materials characterization must be thorough and complete. Design also plays an important role. For example, if a material is selected for a design using mechanical fastening, and nonmechanical joining is substituted as the manufacturing technique, the material and the design must be reevaluated. A design requiring mechanical fastening generally is not optimal for nonmechanical joining.

FIG. 8 INTERRELATIONSHIPS AMONG AEROSPACE DESIGN, MANUFACTURE, MAINTENANCE AND REPAIR, AND MATERIAL SELECTION. SOURCE: REF 8

Thermal Stability. A clear understanding of the material properties of importance for a given application is crucial to optimal material selection. For example, thermal stability, both metallurgical and physical, is a very important factor in elevated-temperature applications. Over time, microstructural thermal instability may alter mechanical properties, such as toughness, creep, or crack growth. The thermal stability of a material may be affected by a thermal joining process. Precipitation of unstable phases in the fusion zones of structural alloys, such as advanced titanium alloys, can be difficult to evaluate. Also, the tensile strength of high-strength aluminum alloys can be adversely affected during high-temperature

curing of adhesives if the aging temperature is exceeded for a sufficient length of time (Ref 9). Thermal stability can be maintained by ensuring proper material chemistry, heat treatment, joining process, joining consumable, and/or postjoining thermal treatments. Thermal instability may manifest itself as creep, cracking, buckling, and/or a reduction in corrosion resistance (Ref 4, 10, 11, 12). The corrosion properties of stainless steels may be severely impaired by elevatedtemperature exposure, such as welding and brazing. The degradation referred to as sensitization can also reduce fatigue performance. Thermal stability has been used here as an example. A similar discussion could be presented for each property listed in Table 13. Weight is an overriding concern in aerospace applications because of its adverse effect on performance and operational

costs. Many decisions are based on the specific property of a material, such as those shown in Fig. 9. Specific properties are defined as the mechanical or physical property of the material divided by its density: Material property = specific property Material density

When using specific properties, one must be careful to fully characterize anisotropic materials. When considering a material for a nonmechanical joining application, an understanding of the effect of joining on these properties is essential. Joinability often is ignored during the development of aerospace materials (and materials in general) and can considerably inhibit their versatility during construction. Because compromise is the standard when selecting materials for a specific application, the effect of joining on properties must be considered.

FIG. 9 AVERAGE VERSUS SPECIFIC PROPERTIES

Specifications. Aerospace materials are most often purchased to a military (MIL) or Aerospace Materials Specification (AMS). Some materials, such as A286, are available in standard and weldable grades. The weldable grade is formulated to reduce the tendency for microcracking. Filler metal specifications are provided by AWS or by the Society of Automotive Engineers (SAE) as Aerospace Materials Specifications. Economics. Cost is another important material consideration. The cost of the raw materials is usually a small percentage

of the cost of the end product in applications such as aerospace airframes and engine components, structures, and assemblies. The costs of some common materials as a function of volume production are presented in Fig. 10. However, material selection can have a significant cost impact in the areas of manufacturability and life-cycle costs (repairability and maintainability). The designation of materials must be as specific as possible to ensure maintenance of quality for joining applications (Ref 10, 11, 13). If multiple welds are to be performed on a given structure, the requirement for intermediate heat treatments or other processing steps must be considered. In large structures where full heat treatments are difficult because of distortion, the reduction in properties may require increased joint thickness, thereby increasing weight and complicating the manufacturing process.

FIG. 10 COST VERSUS PRODUCTION FOR SEVERAL COMMON AEROSPACE MATERIALS. SOURCE: REF 8

Joining process selection and practice are influenced by material selection. For example, titanium requires vacuum or complete inert shielding (chambers with glove ports) to produce high-quality joints when elevated-temperature joining processes are used, which increases production costs in comparison with some other materials. Finally, the environmental ramifications of joining must be considered when selecting a material. Toxic materials, such as beryllium, are sometimes used in aerospace structures because of their very high specific properties. However, when joining these materials, especially by welding, health and safety issues may drive manufacturing costs very high.


4. S. YUKAWA, GUIDELINE FOR PRESSURE VESSEL SAFETY ASSESSMENT, NIST SPEC. PUBL. 780, NATIONAL INSTITUTE FOR STANDARDS AND TECHNOLOGY, P 14-22 7. G.E. DIETER, ENGINEERING DESIGN, MCGRAW-HILL, 1983 8. B.D. TAPLEY, ESCHBACK'S HANDBOOK OF ENGINEERING FUNDAMENTALS, JOHN WILEY & SONS, 1990 9. AEROSPACE STRUCTURAL MATERIALS HANDBOOK, VOL III, 1984 10. SIMS AND HAGEL, THE SUPERALLOYS, JOHN WILEY & SONS, 1972 11. CASTRO AND DECADENT, WELDING METALLURGY OF STAINLESS AND HEAT RESISTING STEELS, CAMBRIDGE PRESS, 1968 12. EVANS AND WILSHIRE, CREEP IN METALS AND ALLOYS, DOTESIOS PRINTERS, U.K., 1985 13. AEROSPACE STRUCTURAL MATERIALS HANDBOOK, VOL IV, 1984 Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Joinability Material Characteristics Related to Joining. Materials for aerospace use can be divided into two major categories: structural and engine, the latter being materials for elevated temperatures. There is considerable overlap, but it is instructive to examine materials from these points of view. Although the practice of welding of critical structural members in modern aircraft is small, but growing, the use of welding in missiles and space hardware has always been extensive. The joinability of aerospace materials takes on various definitions. Joint efficiency, cracking resistance, welding characteristics (such as flow), maintenance of corrosion resistance, and long-term metallurgical stability are all measures of joinability. Design, process, and consumable selection affect the joining characteristics of materials.

Dissimilar-materials joining muddies these definitions even further. For example, the welding of titanium to stainless steel is not recommended, but brazing these two materials together is possible. The usefulness of the hybrid braze would be determined by the service requirements. Joinability is often determined by standardized tests or by the ability of a material to meet certain mechanical or inspection requirements. When evaluating joinability, it is important to consider the range of conditions to be experienced n production. For example, most material specifications allow a range of compositions. Because joinability may vary considerably as a function of chemistry, addressing these concerns prior to production will avoid costly rework. The joinability of a material is influenced by the condition of the material when joined. For example, nickel-base superalloys should always be stress relieved prior to brazing to prevent stress cracking (Ref 14). Heat-treatable aluminum alloys are sometimes welded in the T4 condition and then postweld aged to the T6 condition to achieve more uniform properties in the weldment. Alpha-beta titanium alloys are sensitive to preweld heat treatment. The thermal history can become very important during repair operations. If a material has been exposed to elevated temperatures or high loading in service, its joinability may be altered. In such instances, specialized testing may be required to ensure adequate joinability for repair operation. The joinability of a material may be contingent on the postjoining operations, such as heat treatments. Many as-joined materials do not have the properties required for a particular application, but can be brought to an acceptable condition by a postjoining process. Postjoining thermal treatments can consist of full heat treatments (solution, intermediate, age), stress relief, or stabilization treatments (age). These can be localized or can be applied to the entire part. Other postjoining treatments include shot peening, machining, vibration stress relief, and sizing. Base Materials. When characterizing the joinability of a material, it is important to consider the effect of joining on the

base material as well as on the nexus (the actual joint; i.e., the fusion zone in welding). Nonthermal joining operations, such as room-temperature-cured adhesive bonding, may have no effect on the base material. In contrast, thermal

processes, such as brazing, soldering, and welding, can radically change material properties, especially near the nexus. Characterization of the affected base material can be difficult because of its size and makeup. For example, the HAZ in an electron beam weld in thick-section titanium may be as little as half the fusion-zone width. If the weld is on the order of 2.5 mm (0.10 in.) in width, the HAZ is only 1.25 mm (0.05 in.) in width. Locating a notch in the HAZ to evaluate fatigue or fracture toughness can be challenging. Moreover, if the HAZ is small enough, its effect on the service performance of the joint may be negligible, which signals a caution when service life is evaluated by using "simulated" HAZ microstructure. Joints of the same configuration as the production parts should be used whenever possible. Filler Materials. The selection of filler materials is a critical factor that sometimes receives less than proper attention when joining aerospace materials (Ref 15). Filler materials are most often selected on the basis of chemistry match. However, other considerations sometimes take precedence, such as ferrite control in the welding of stainless steels, distortion/residual stress control for high-strength steels, optimization of elevated-temperature properties for nickel-base superalloys, and strength optimization in aluminum welding or optimization of environmental resistance in brazing and adhesive bonding.

Selection of filler materials for dissimilar-material joining requires careful consideration. Some brazing alloys contain melting-point depressants to allow brazing at temperatures lower than the operating environment (diffusion brazing). The effect of the formation of detrimental phases in the base materials is a factor when these alloys are used. Special corrosion or wear requirements will sometimes direct the use of a particular joining consumable. Joining Techniques Other Than Fusion Welding. Process is another consideration in the joinability of aerospace

materials. Alloys that are nonfusion weldable may be readily joined by other methods, such as friction, explosion, diffusion, brazing, or soldering. Service Considerations. The joinability of materials is influenced by service requirements. A titanium weldment used

for a missile fuselage or a wing support must have high strength, low fatigue crack-growth rates, and high toughness. Good corrosion resistance and thermal stability are also desirable. The properties of the weldments made from such material must reflect the base material requirements. On the other hand, an Inconel 626 weld used to build up a turbine blade tip may not have the same requirements as the turbine blade. Ease of welding and adequate elevated-temperature oxidation resistance are overriding concerns in this case. Because the tip buildup is not "structural" in nature and the consequences of failure not as severe as for the wing support, weldability is based on different criteria. Therefore, the joinability of a given material can vary depending on the end service requirements.


14. BRAZING HANDBOOK, AMERICAN WELDING SOCIETY, 1976 15. D. HARVEY, FILLER METAL SELECTIONS FOR CRITICAL AEROSPACE FABRICATION, WELD. DESIGN FABR., VOL 60 (NO. 2), MARCH 1991, P 765-80 Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Damage Tolerance Existence of Flaws. Damage tolerance is a design methodology often used for aerospace systems. The basic premise is that all structures contain inherent flaws. The effect of these flaws on flight safety depends on their geometry, growth rate at the service conditions, and critical size, and on the inspectability of the structure and the consequence of failure (Ref 16).

In the design of damage-tolerant structures, it is imperative that the relevant material properties, such as fatigue crackgrowth rates, fracture toughness, and SCC, be well characterized. When applying damage-tolerant methodologies to structures joined by other methods than mechanical, unique considerations include:

• • • • • •

RESIDUAL STRESS EFFECTS PROPERTY GRADIENTS ACROSS THE JOINT INHERENT FLAWS, SUCH AS MICROFISSURES AND POROSITY LONG-TERM STABILITY CORROSION STRESS CONCENTRATIONS

These concerns are not well characterized for many of the materials and joining processes used in aerospace. Finally, because the properties of the base metal may be altered nonuniformly during thermal joining, specialized testing may be required. For example, the crack-growth rate may be improved in the fusion zone of a weld, but the fracture toughness and SCC resistance may be reduced in the HAZ. All of these factors are important to the damage-tolerance analysis, and thus full characterization is required. Inspection Criteria. For complex structures, the limiting factor may well be the inspectability of the joint. As a result,

many welding specifications used for aerospace production, such as MIL-STD-2219 and AMS 2680, include defect allowable tables. When contractually requiring one of these standards, it is important to review the defect allowables to ensure that they are realistic for the application. Many unnecessary repairs are performed on welded structures because of improper defect criteria. As a result of the additional thermal cycles, the repairs may be more detrimental than the flaw in many materials. For many joining processes and materials combinations used in aerospace, the information necessary to perform a damage-tolerance analysis is not available; therefore, the defect criteria in the welding specifications are used, as this procedure is less expensive than developing the required data. Also, the effect of residual stress on fracture toughness and fatigue crack growth is difficult to evaluate with current damage-tolerance methodologies. Much work needs to be accomplished in this area. Nondestructive examination (NDE) concerns require consideration of both the process and design limitations. Many times joints are designed without adequate attention to inspection requirements, often making effective NDE impossible due to inaccessibility.


16. USAF DAMAGE TOLERANT DESIGN HANDBOOK: GUIDELINES FOR ANALYSIS AND DESIGN OF DAMAGE TOLERANT AIRCRAFT, U.S. AIR FORCE, 1984 Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Advanced Materials Some materials used for aerospace systems present special problems in relation to joining. Among these materials are composites and intermetallics. Composite materials, including metal matrix and organic matrix, are used because of their increased stiffness, damping properties, and high specific properties. Composites present challenges when joined by any method, and much work is currently underway to establish their joinability. Maintenance of properties when using nonmechanical joining is difficult. Metal-matrix composites have been successfully joined by brazing, diffusion, and high-energy density welding. During the development of these materials, joinability was not a high priority and is thus limited. The bond integrity between the fibers and the matrix of a composite greatly affects material properties. Special coatings are applied to the fibers to ensure adequate bonding. The stability of these coatings to a large degree determines joinability when using elevated-temperature processes. Coating schemes that optimize joinability should be addressed during materials development. A further challenge involves the attainment of material properties across the joint.

Intermetallics, such as Ti3Al2, offer very promising elevated-temperature properties. Limited studies to date concerning the fusion weldability of these materials indicate that tight control of base-metal chemistry, accurate regulation of heat input and cooling rate, and joint design to minimize residual stress are required to produce sound joints (Ref 17). Brazing offers some advantages, but may be limited by operating temperature unless suitable diffusion-bonding fillers can be developed. The joinability of these types of emerging materials will to a large degree determine their long-term usefulness in aerospace systems.


17. PATTERSON ET AL. TITANIUM ALUMINIDE: ELECTRON BEAM WELDABILITY," WELD. J., JAN 1990, P 39S-44S Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Machinery and Equipment

Welding plays an important role in the wide variety of structures broadly characterized as machinery and equipment. This section will discuss three categories: stationary equipment, such as presses, rolling mills, and machine tools; mobile equipment, such as cranes, earthmoving, and other construction equipment; and rotating components of machines. In all of these categories, the materials requirements for serviceability are distinctly different for the structural machine support parts and for those involving the machine-moving and working parts. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Applications Stationary Equipment. The working parts of machine tools require stability. Massive cast iron or cast steel bases are used to provide a damping capacity, thus ensuring that cutting tools do not vibrate or chatter. In applications that use weldments, thick steel components are specified.

The principal design requirement for metalworking presses is adequate strength to counter the forces imposed during press operation. Weldments of various sections using both wrought and cast carbon steels offer adequate strength and rigidity. The elastic deflections of the components created by work forces affect the design, which must take into account the cyclical nature of the stress and hence the fatigue properties of the materials and welded joints. Mobile Equipment. In the design of components for earthmoving and other construction equipment, static and dynamic

loads are important factors, as are erosion, wear, weight, and ease of maintenance. The materials for each working part of such equipment must be considered separately. As an example, booms used in mobile industrial cranes must be light in weight, yet rigid under load. Wear resistance is another important factor in this application. Rotating Elements. Materials specified for the rotating parts of pumps, fans, compressors, and similar machinery have a

wide variety of properties. For large rotating components, weldments are commonly required, sometimes employing different metals for improved performance in specific environments. Erosion, corrosion, and cavitation are often the

mechanisms that limit the lives of such machine elements. Critical surfaces are frequently weld surfaced with corrosionresistant materials. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Design Considerations Welds and members of weldments are sized to withstand the working stress and stress range induced by static and fluctuating loads. The magnitude of the fluctuation between maximum and minimum working stress is termed the working stress range. The magnitude of the static and fluctuating stresses is influenced by such factors as dead load, live load, type of operating mechanism, machine geometry, component size, shock or impact, accidental overloads, corrosion, wear misalignment, stress concentration, and weld joint design. The allowable stress and stress range must be greater than the expected working stress and stress range. The relationships among the various stresses, stress ranges, strengths, and loads are shown graphically in Fig. 11.

FIG. 11 STRESS RELATIONSHIPS RESULTING FROM STATIC AND FLUCTUATING LOADS

Allowable stresses in weld and base metals are usually specified by an applicable code or standard based on a history of satisfactory operating performance. In the absence of an industry standard, as in the case of most machinery and equipment, the determination of an allowable stress value is left to the discretion of the manufacturer or designer. In such case, the allowable stress (SA) is calculated by dividing the ultimate tensile strength (UTS) or yield strength (YS) by an appropriate factor of safety (FS) selected by the designer: SA =

UTS YS or FS FS

The selection of an appropriate factor of safety to prevent failure from unpredictable causes is empirical and depends on industry data and the experience and judgment of the designer. Table 14 provides guidance in selecting a factor of safety, based on the ultimate tensile strength of the material. The assigned factors in Table 14 may be increased as warranted by safety, cost, plant utilization, and engineering judgment. Also, the assigned factors may be reduced if detailed analysis increases confidence in design factors, such as loading, material properties, and fabrication controls.

TABLE 14 FACTOR OF SAFETY SELECTION CRITERIA

CONDITION

EXCEPTIONALLY RELIABLE MATERIALS USED UNDER CONTROLLABLE CONDITIONS AND SUBJECTED TO LOADS AND STRESSES THAT CAN BE DETERMINED WITH CERTAINTY. USED ALMOST INVARIABLY WHERE LOW WEIGHT IS A PARTICULARLY IMPORTANT CONSIDERATION WELL-KNOWN MATERIALS, UNDER REASONABLY CONSTANT ENVIRONMENTAL CONDITIONS, SUBJECTED TO LOADS AND STRESSES THAT CAN BE DETERMINED LESS TRIED OR BRITTLE MATERIALS UNDER AVERAGE CONDITIONS OF ENVIRONMENT, LOAD, AND STRESS KNOWN MATERIALS THAT ARE TO BE USED IN UNCERTAIN ENVIRONMENTS OR SUBJECTED TO UNCERTAIN STRESSES

ASSIGNED FACTOR OF SAFETY(A) 2.5-3

4-5

5-6 6-8

(A) FOR REPEATED LOADS, THE FACTORS SHOULD BE DOUBLED. FOR IMPACT LOADS. THE FACTORS SHOULD BE INCREASED AS WARRANTED. When a design is analyzed using a mathematical model, it must be understood that subsequent calculations only approximate the working stress. With all models, various assumptions, limitations, and simplifications used to fit equations to the component model influence the accuracy of the calculations. (This is true even when using finite-element methods.) The calculated stress values that result from applying "good" theory to a "good" model are, at best, ballpark figures that give the design engineer a rational basis for making design decisions. For the most part, the allowable stress ranges specified in current standards for welded joints subjected to fatigue loading are based on testing of representative full-size welded joints in actual or mockup structures. Static Loads. Nearly every machine is subjected to static and fluctuating loads; few useful machine structures have only

static loads applied. Structures subjected to slowly applied loads, few cycles (less than 20,000) of loading, or nearly steady loads are considered statically loaded and are designed using rules that apply to static structures. These rules assume that localized areas of high stress due to stress concentrations are relieved by yielding of the highly stressed material to redistribute internal stresses. Steels commonly used for welded construction of machinery and equipment have yield strengths of less than 700 MPa (100 ksi) and are listed in Table 15. Class I materials have the highest weldability, and class V have the lowest. Carbon and carbon-manganese steels--those class I and class II steels designated by the American Iron and Steel Institute (AISI), SAE, and ASTM--are the most commonly used materials for welded machinery and equipment. These steels do not require preheat unless the thickness is greater than 150 mm (1

1 in.). High-strength low-alloy steels (classes III, IV, and 2

V) allow the thickness and hence the weight of a machine structure subjected to high static loads to be reduced.

TABLE 15 WELDABILITY CLASSES OF STEELS USED IN MACHINERY AND EQUIPMENT

STEEL

TYPICAL MINIMUM YIELD

CARBON SPECIFICATION EQUIVALENT (CEQ)

STRENGTH (MAX)(A) MPA KSI STRUCTURAL CARBON 20030-60 0.38 CLASS 400 I

CLASS II

AISI-SAE

ASTM

API

...

...

A 36, A 38 (GRADE Y35), A 53 (GRADE B), A 106, A 131, A 139, A 374, A 500, A 501, A 516, A 524, A 529, A 570 (GRADES D, E), A 573

52 (GRADE B)

5LX (GRADE 42)

...

A 242, A 441, A 537 (GRADES A, E), A 572 (GRADES 42, 45, 50), A 588, A 618 A 572 (GRADES 60, 65)

...

A 514, A 517

...

240380

35-55 0.48

1005, 1006, 1008, 1010, 1012, 1015, 1016, 1017, 1018, 1020, 1021 ...

275380

40-55 0.63

...

HSLA CLASS III

41060-65 NOT CLASS 450 SPECIFIED IV QUENCHED AND TEMPERED 620900.74 CLASS 700 100 V

...

(A) CARBON EQUIVALENT IS A MEASURE OF WELDABILITY. FOR UNALLOYED (STRUCTURAL) STEELS. CEQ = %C + (MN/4) + (SI/4). FOR LOW-ALLOY STEELS, CEQ = %C + (MN/6) + (NI/20) + (CR/10) + (MO/40) + (V/10). The AISI and SAE carbon-manganese steels have carbon contents of less than 0.25% and manganese contents of less than 0.90% and are supplied to meet composition specifications that do not include mechanical properties. Unless verified by mechanical tests, an ultimate tensile strength of 380 MPa (55 ksi) and a yield strength of 190 MPa (27.5 ksi) should be used for design purposes. These steels have the highest weldability rating shown in the machinery welding specifications that classify steels according to weldability. ASTM material specifications include mechanical property requirements that demand higher strength than the AISI and SAE carbon-manganese steels, and thus offer the opportunity to reduce thickness and weight for statically loaded conditions. The weldability of these higher-strength steels is lower, and they usually require higher preheat and lowhydrogen welding processes. Allowable stress for statically loaded steel base metals can be determined from the equation given above after selecting a factor of safety. Allowable stress levels for statically loaded weld metal are presented in Table 16.

TABLE 16 ALLOWABLE STRESSES IN WELD METAL

TYPE OF WELD

STRESS IN WELD

ALLOWABLE STRESS

GROOVE WELDS WITH COMPLETE JOINT PENETRATION

TENSION NORMAL TO EFFECTIVE AREA COMPRESSION NORMAL TO

SAME AS BASE METAL

SAME AS BASE METAL

REQUIRED WELD STRENGTH LEVEL(A) MATCHING WELD METAL SHALL BE USED. WELD METAL WITH A STRENGTH LEVEL EQUAL

EFFECTIVE AREA

TENSION OR COMPRESSION PARALLEL TO WELD AXIS SHEAR ON EFFECTIVE AREA

GROOVE WELDS WITH PARTIAL JOINT PENETRATION

COMPRESSION NORMAL TO EFFECTIVE AREA: JOINT NOT DESIGNED TO BEAR

COMPRESSION NORMAL TO EFFECTIVE AREA; JOINT DESIGNED TO BEAR TENSION OR COMPRESSION PARALLEL TO WELD AXIS(B) SHEAR PARALLEL TO WELD AXIS

TENSION NORMAL TO EFFECTIVE AREA

FILLET WELDS

SHEAR ON EFFECTIVE AREA

SAME AS BASE METAL

0.27 × NOMINAL TENSILE STRENGTH OF WELD METAL, EXCEPT SHEAR STRESS ON BASE METAL SHALL NOT EXCEED 0.40 × YIELD STRENGTH OF BASE METAL. 0.45 × NOMINAL TENSILE STRENGTH OF WELD METAL, EXCEPT STRESS ON BASE METAL SHALL NOT EXCEED 0.55 × YIELD STRENGTH OF BASE METAL. SAME AS BASE METAL

TO OR ONE CLASSIFICATION (69 MPA, OR 10 KSI) LESS THAN MATCHING WELD METAL MAY BE USED. WELD METAL WITH A STRENGTH LEVEL EQUAL TO OR LESS THAN MATCHING WELD METAL MAY BE USED.

WELD METAL WITH A STRENGTH LEVEL EQUAL TO OR LESS THAN MATCHING WELD METAL MAY BE USED.

SAME AS BASE METAL

0.27 × NOMINAL TENSILE STRENGTH OF WELD METAL. EXCEPT SHEAR STRESS ON BASE METAL SHALL NOT EXCEED 0.36 × YIELD STRENGTH OF BASE METAL. 0.27 × NOMINAL TENSILE STRENGTH OF WELD METAL, EXCEPT SHEAR STRESS ON BASE METAL SHALL NOT EXCEED 0.55 × YIELD STRENGTH OF BASE METAL. 0.27 × NOMINAL TENSILE STRENGTH OF WELD METAL

WELD METAL WITH A STRENGTH LEVEL EQUAL TO OR LESS THAN

PLUG AND SLOT WELDS

TENSION OR COMPRESSION PARALLEL TO WELD AXIS(B) SHEAR PARALLEL TO FAYING SURFACES (ON EFFECTIVE AREA)

SAME AS BASE METAL

MATCHING WELD METAL MAY BE USED

0.27 × NOMINAL TENSILE STRENGTH OF WELD METAL, EXCEPT SHEAR STRESS ON BASE METAL SHALL NOT EXCEED 0.36 × YIELD STRENGTH OF BASE METAL.

WELD METAL WITH A STRENGTH LEVEL EQUAL TO OR LESS THAN MATCHING WELD METAL MAY BE USED.

(A) FOR MATCHING WELD METAL, SEE TABLE 20. (B) FILLET WELDS AND GROOVE WELDS WITH PARTIAL JOINT PENETRATION JOINING THE COMPONENT ELEMENTS OF BUILT-UP MEMBERS, SUCH AS FLANGE-TO-WEB CONNECTIONS, MAY BE DESIGNED WITHOUT REGARD TO THE TENSILE OR COMPRESSIVE STRESS IN THESE ELEMENTS PARALLEL TO THE AXIS OF THE WELDS. Fluctuating Loads. Fatigue is the process of progressive localized damage that may result in cracks or complete fracture after a sufficient number of working stress fluctuations. Fatigue life is the number of stress fluctuations sustained to failure.

The fatigue design of welded joints is different from the "modified Goodman diagram" method used for plain (unwelded) machine and structural members. The fatigue life of plain base metal is a function of its tensile strength. Welded joints introduce stress concentrations in a structure, which reduces fatigue life. Examples of stress concentrations include weld ripple, undercut, crowned reinforcement, lack of penetration, slag inclusions, microcracks, and craters. The fatigue life of a steel welded joint is independent of base metal tensile strength. An ASTM A 36 steel weld joint has the same fatigue strength as a joint made from ASTM A 514 steel. The fatigue life of a steel weld joint is indirectly proportional to working stress range; that is, as the magnitude of working stress range increases, fatigue life decreases (see Table 17). The fatigue life of a steel weld joint is a function of the weld joint category, which in turn is a function of joint geometry, type of weld, direction of load application, and weld quality assurance (see Table 18 and Fig. 12).

TABLE 17 ALLOWABLE STRESS RANGES

NOMINAL NUMBER OF LOADING CYCLES STRESS CATEGORY 20,000100,000500,000>2,000,000 (SEE TABLE 18) 100,000 500,000 2,000,000 MPA KSI MPA KSI MPA KSI MPA KSI A 434 63 255 37 165 24 165 24 B 338 49 200 29 124 18 110 16 B' 269 39 159 23 103 15 83 12 C 241 35 145 21 90 13 69(A) 10(A) D 193 28 110 16 69 10 48 7 E 152 22 90 13 55 8 31 4.5 E' 110 16 63 9.2 40 5.8 18 2.6 F 103 15 83 12 62 9 55 8 (A) FLEXURAL STRESS RANGE OF 83 MPA (12 KSI) PERMITTED AT TOE OF STIFFENER WELDS OR FLANGES TABLE 18 STRESS CATEGORY CLASSIFICATIONS

GENERAL CONDITION

SITUATION

TYPES OF STRESS(A)

STRESS CATEGORY (SEE TABLE 17)

ILLUSTRATIVE EXAMPLE NOS. (SEE FIG. 12)(B)

PLAIN MATERIAL

BASE METAL WITH ROLLED OR CLEANED SURFACE. FLAME-CUT EDGES WITH ANSI SMOOTHNESS OF 1000 OR LESS BASE METAL IN MEMBERS WITHOUT ATTACHMENTS, BUILT-UP PLATES, OR SHAPES CONNECTED BY CONTINUOUS FULL-PENETRATION GROOVE WELDS OR BY CONTINUOUS FILLET WELDS PARALLEL TO THE DIRECTION OF APPLIED STRESS BASE METAL IN MEMBERS WITHOUT ATTACHMENTS, BUILD-UP PLATES, OR SHAPES CONNECTED BY FULLPENETRATION GROOVE WELDS WITH BACKING BARS NOT REMOVED, OR BY PARTIALPENETRATION GROOVE WELDS PARALLEL TO THE DIRECTION OF APPLIED STRESS BASE METAL AT TOE WELDS ON GIRDER WEBS OR FLANGES ADJACENT TO WELDED TRANSVERSE STIFFENERS BASE METAL AT ENDS OF PARTIALLENGTH WELDED COVER PLATES NARROWER THAN THE FLANGE, HAVING SQUARE OR TAPERED ENDS, WITH OR WITHOUT WELDS ACROSS THE ENDS OR WIDER THAN FLANGE WITH WELDS ACROSS THE ENDS FLANGE THICKNESS ≤ 20 MM (0.8 IN.) FLANGE THICKNESS >20 MM (0.8 IN.) BASE METAL AT END OF PARTIALLENGTH WELDED COVER PLATES WIDER THAN THE FLANGE WITHOUT WELDS ACROSS THE ENDS BASE METAL AND WELD METAL AT FULL-PENETRATION GROOVE WELDED SPLICES OF PARTS OF SIMILAR CROSS SECTION GROUND FLUSH, WITH GRINDING IN THE DIRECTION OF APPLIED STRESS AND WITH WELD SOUNDNESS ESTABLISHED BY RADIOGRAPHIC OR ULTRASONIC INSPECTION IN

T OR REV

A

1, 2

T OR REV

B

3-6

T OR REV

B'

3-6

T OR REV

C

7

T OR REV T OR REV

E

5

E'

5

E'

5

B

10, 11

BUILT-UP MEMBERS

GROOVE WELDS

T OR REV

ACCORDANCE WITH THE REQUIREMENTS OF 9.25.2 OR 9.25.3 OF AWS D1. I BASE METAL AND WELD METAL AT FULL-PENETRATION GROOVE WELDED SPLICES AT TRANSITIONS IN WIDTH OR THICKNESS; WITH WELDS GROUND TO PROVIDE 1 2

SLOPES NO STEEPER THAN 1 TO 2 , WITH GRINDING IN THE DIRECTION OF APPLIED STRESS AND WITH WELD SOUNDNESS ESTABLISHED BY RADIOGRAPHIC OR ULTRASONIC INSPECTION IN ACCORDANCE WITH THE REQUIREMENTS OF 9.25.2 OR 9.25.3 OF AWS D1. 1 A 514 BASE METAL OTHER BASE METALS BASE METAL AND WELD METAL AT FULL-PENETRATION GROOVE WELDED SPLICES, WITH OR WITHOUT TRANSITIONS HAVING SLOPES NO GREATER THAN 1 TO 2

PARTIALPENETRATION GROOVE WELDS

FILLET-WELDED CONNECTIONS

BASE METAL AT MEMBERS

B'

12, 13

B

12, 13

C

10-13

T OR REV

F(C)

16

T OR REV

E

...

T OR REV T OR REV

E

17, 18

E'

17, 18

1 2

WHEN REINFORCEMENT IS NOT REMOVED BUT WELD SOUNDNESS IS ESTABLISHED BY RADIOGRAPHIC OR ULTRASONIC INSPECTION IN ACCORDANCE WITH THE REQUIREMENTS OF 9.25.2 OR 9.25.3 OF AWS D1.1 WELD METAL OF PARTIALPENETRATION TRANSVERSE GROOVE WELDS, BASED ON EFFECTIVE THROAT AREA OF THE WELD OR WELDS BASE METAL AT INTERMITTENT FILLET WELDS BASE METAL AT JUNCTION OF AXIALLY LOADED MEMBERS WITH FILLET-WELDED END CONNECTIONS. WELDS SHALL BE DISPOSED ABOUT THE AXIS OF THE MEMBER SO AS TO BALANCE WELD STRESSES. B ≤ 25 MM (1 IN.) B >25 MM (1 IN.)

T OR REV T OR REV T OR REV

CONNECTED WITH TRANSVERSE FILLET WELDS B ≤ 13 MM (0.5 IN.)

FILLET WELDS

PLUG OR SLOT WELDS MECHANICALLY FASTENED CONNECTIONS

ATTACHMENTS

B > 13 MM (0.5 IN.) WELD METAL OF CONTINUOUS OR INTERMITTENT LONGITUDINAL OR TRANSVERSE FILLET WELDS BASE METAL AT PLUG OR SLOT WELDS SHEAR ON PLUG OR SLOT WELDS BASE METAL AT GROSS SECTION OF HIGH-STRENGTH BOLTED SLIPCRITICAL CONNECTIONS, EXCEPT AXIALLY LOADED JOINTS THAT INDUCE OUT-OF-PLANE BENDING IN CONNECTED MATERIAL BASE METAL AT NET SECTION OF OTHER MECHANICALLY FASTENED JOINTS BASE METAL AT NET SECTION OF FULLY TENSIONED HIGHSTRENGTH, BOLTED-BEARING CONNECTIONS BASE METAL AT DETAILS ATTACHED BY FULL-PENETRATION GROOVE WELDS SUBJECT TO LONGITUDINAL AND/OR TRANSVERSE LOADING WHEN THE DETAIL EMBODIES A TRANSITION RADIUS, R, WITH THE WELD TERMINATION GROUND SMOOTH AND FOR TRANSVERSE LOADING, THE WELD SOUNDNESS ESTABLISHED BY RADIOGRAPHIC OR ULTRASONIC INSPECTION IN ACCORDANCE WITH 9.25.2 OR 9.25.3 OF AWS D1.1 LONGITUDINAL LOADING R > 600 MM (24 IN.) 600 MM (24 IN.) > R > 150 MM (6 IN.) 150 MM (6 IN.) > R > 50 MM (2 IN.) 50 MM (2 IN.) > R DETAIL BASE METAL FOR TRANSVERSE LOADING: EQUAL THICKNESS AND REINFORCEMENT REMOVED R > 600 MM (24 IN.)

T OR REV

C(C)

20, 21

S

F(C)

15, 17, 18, 20, 21

T OR REV S T OR REV

E

27

F B

27 8

T OR REV

D

8, 9

T OR REV

B

8, 9

T OR REV T OR REV T OR REV T OR REV

B

14

C

14

D

14

E

14

T OR REV

B

14

600 MM (24 IN.) > R > 150 MM (6 IN.) 150 MM (6 IN.) > R > 50 MM (2 IN.) 50 MM (2 IN.) > R DETAIL BASE METAL FOR TRANSVERSE LOADING; EQUAL THICKNESS AND REINFORCEMENT NOT REMOVED R > 600 MM (24 IN.) 600 MM (24 IN.) > R > 150 MM (6 IN.) 150 MM (6 IN.) > R > 50 MM (2 IN.) 50 MM (2 IN.) > R DETAIL BASE METAL FOR TRANSVERSE LOADING: UNEQUAL THICKNESS AND REINFORCEMENT REMOVED R > 50 MM (2 IN.) 50 MM (2 IN.) > R DETAIL BASE METAL FOR TRANSVERSE LOADING; UNEQUAL THICKNESS AND REINFORCEMENT NOT REMOVED ALL R DETAIL BASE METAL FOR TRANSVERSE LOADING R > 150 MM (6 IN.) 150 MM (6 IN.) > R > 50 MM (2 IN.) 50 MM (2 IN.) > R


C

14

D

14

E

14, 15


C

14

C

14

D

14

E

14, 15

T OR REV

D

14

E

14, 15

T OR REV

E

14, 15


C

19

D

19

E

19

T OR REV

D

15

T OR REV T OR REV

E

15

E'

15

BASE METAL AT DETAIL ATTACHED BY FULL-PENETRATION GROOVE WELDS SUBJECT TO LONGITUDINAL LOADING 50 MM (2 IN.) < A < 12B OR 100 MM (4 IN.) A > 12B OR 100 MM (4 IN.) WHEN B ≤ 25 MM (1 IN.) A > 12B OR 100 MM (4 IN.) WHEN B > 25 MM (1 IN.) BASE METAL AT DETAIL ATTACHED BY FILLET WELDS OR PARTIALPENETRATION GROOVE WELDS

SUBJECT TO LONGITUDINAL LOADING A < 50 MM (2 IN.) 50 MM (2 IN.) < A < 12B OR 100 MM (4 IN.) A > 12B OR 100 MM (4 IN.) WHEN B ≤ 25 MM (1 IN.) A > 12B OR 100 MM (4 IN.), WHEN B > 25 MM (1 IN.) BASE METAL ATTACHED BY FILLET WELDS OR PARTIAL-PENETRATION GROOVE WELDS SUBJECTED TO LONGITUDINAL LOADING WHEN THE WELD TERMINATION EMBODIES A TRANSITION RADIUS WITH THE WELD TERMINATION GROUND SMOOTH R > 50 MM (2 IN.) R ≤ 50 MM (2 IN.) FILLET-WELDED ATTACHMENTS WHERE THE WELD TERMINATION EMBODIES A TRANSITION RADIUS, WELD TERMINATION GROUND SMOOTH, AND MAIN MATERIAL SUBJECT TO LONGITUDINAL LOADING DETAIL BASE METAL FOR TRANSVERSE LOADING: R > 50 MM (2 IN.) R < 50 MM (2 IN.) BASE METAL AT STUD-TYPE SHEAR CONNECTOR ATTACHED BY FILLET WELD OR AUTOMATIC END WELD SHEAR STRESS ON NOMINAL AREA OF STUD-TYPE SHEAR CONNECTORS


C

15, 23-26

D

15, 23, 24, 26

E

15, 23, 24, 26

E'

15, 23, 24, 26

T OR REV T OR REV

D

19

E

19


D

19

E

19

C

22

S

F

...

(A) T, RANGE IN TENSILE STRESS ONLY; REV, RANGE INVOLVING REVERSAL OF TENSILE OR COMPRESSIVE STRESS; S, RANGE IN SHEAR, INCLUDING SHEAR-STRESS REVERSAL. (B) THESE EXAMPLES ARE PROVIDED AS GUIDELINES AND ARE NOT INTENDED TO EXCLUDE OTHER REASONABLY SIMILAR SITUATIONS. (C) ALLOWABLE FATIGUE STRESS RANGE FOR TRANSVERSE PARTIAL-PENETRATION AND TRANSVERSE FILLET WELDS IS A FUNCTION OF THE EFFECTIVE THROAT, DEPTH OF PENETRATION, AND PLATE THICKNESS.

FIG. 12 ILLUSTRATIVE EXAMPLES OF THE STRESS CATEGORIES IN TABLE 18

The primary design requirement for some machine members can be rigidity. Such members, even when subjected to both static and fluctuating loads, are treated as statically loaded structures. Deflection under load is very small and the working stress is only 14 to 28 MPa (2 to 4 ksi). In this case, the weld need not be sized to carry the full load capacity of the member. Loads are considered cyclic if the stress range exceeds 10% of the allowable stress range or if the total life cycles exceed 20,000. Care must be taken when selecting higher-strength materials, because the allowable stress range of welded joints is independent of material strength. The maximum allowable stress is higher for higher-strength steels, but the allowable stress range is the same. The use of higher-strength steels in cyclic load applications is beneficial only when the part is not welded or when the part is in a tensile prestress condition. The materials listed in Table 15 are commonly used steels that are considered prequalified. Other steels, such as AISI 4340 and 4140, are also commonly used in machinery components, but are not prequalified. Weld procedure qualifications must be developed as required by the appropriate machinery weld specifications. Preheats for these highhardenable steels must be above the martensitic start (Ms) temperature of 290 °C (550 °F) for AISI 4340 and 4140 to allow the formation of the more ductile bainitic structure and thus avoid brittleness and HAZ cracks. A low-hydrogen welding process is mandatory. The postheat temperature must be no higher than 28 °C (50 °F) below the hardening temper temperature incurred when the part was prehardened. A table of welding practices applicable to most steels for machinery and equipment is available in Ref 18. Related AWS specifications are listed in Table 19.

TABLE 19 AWS SPECIFICATIONS RELATING TO MACHINERY AND EQUIPMENT

AWS TITLE NO. D14.1 "WELDING INDUSTRIAL AND MILL CRANES AND OTHER MATERIAL HANDLING EQUIPMENT" D14.2 "METAL CUTTING MACHINE TOOL WELDMENTS" D14.3 "WELDING EARTHMOVING AND CONSTRUCTION EQUIPMENT" D14.4 "CLASSIFICATION AND APPLICATION OF WELDED JOINTS FOR MACHINERY AND EQUIPMENT" D14.5 "WELDING OF PRESSES AND PRESS COMPONENTS" D14.6 "WELDING ROTATING ELEMENTS OF EQUIPMENT" Service Life. Material selection for machine tools, presses, and heavy rolling and forging equipment is usually governed

by an indefinite life expectancy. The working parts (e.g., tools, hammers, forming or punching dies, and mill rolls) and the moving parts (e.g., motors, gears, and clutches) are subject to wear and may need to be either replaced or rebuilt. The basic equipment, however, should be constructed from materials that will last until the machine is retired due to obsolescence. Obsolescence is also a design consideration in material selection for mobile equipment. Because of weight restrictions, however, material election must make use of the maximum load-carrying capacity of the principal components. The life of wear surfaces, especially in earth-moving equipment, is often enhanced by careful selection of weld surfacing materials. Maintenance and repair of components must also factor into the material selection process. If it is known that parts

will need replacement, access for this type of maintenance should be part of the design. If welds must be removed to provide access, necessitating rewelding, the weldability of the steel used for that part is important. A design that seems perfectly feasible during the initial fabrication of a component may present difficult, or even insoluble, problems during disassembly and reassembly for replacement or repair of a worn part. Safety. Construction equipment is used in a variety of applications, some of which may endanger the operator or other

workers at the site. The designer is responsible for ensuring that the material selected is appropriate for the intended use, making allowances for improper uses that cause stress higher than the rated load stress. Material selection should therefore be based on conservative estimates for components whose failure might jeopardize operators and others. Economics. Machinery and equipment involve the same economic principles with regard to material choice as other

types of welded structures. In most applications the unit cost of the material is a small percentage of the overall cost of the fabricated component or machine. Too often the driving force for low initial cost fails to consider the life-cycle cost of a machine. The cost of maintenance, repair, or replacement of components subject to failure or wear can in many cases justify the selection of more costly materials with longer life cycles. Pumps that handle corrosive liquids and dredges that handle erosive slurries are useful examples. Impellers and casings made of materials resistant to corrosion and erosion can add enormously to the utility of a pump over its anticipated life in comparison with materials that may require frequent replacement or rebuilding. In some cases, maintaining the efficiency of a machine component during its life by selecting materials with greater resistance to the environment will offset its cost. Life-cycle cost assessments should be applied to the choice of materials for machinery and equipment during the design phase.


18. R.D. STOUT, WELDABILITY OF STEELS, 4TH ED., WELDING RESEARCH COUNCIL, 1987, P 338437

Material Requirements for Service Conditions

ASM Committee on Material Requirements for Service Conditions*

Material Properties of Weldments Assessment of the performance of welded components must include the ability of the various metallurgical structures to perform the tasks specified by the designer. These structures are broadly defined as the fusion zone and the heat-affected zone. The thermal excursions that occur during welding invariably result in residual stresses, which in some instances may add considerably to the service stresses encountered by a weldment. The effect of welds on the service-ability of weldments for machinery and equipment may influence material selection. Heat-Affected Zone. In unalloyed steels, the properties of a region that has experienced one or more thermal excursions during welding are quite similar to those of the unaffected base metal. In low-alloy steels, especially those in weldability classes IV and V (Table 15), the HAZ becomes considerably stronger, harder, and less ductile. This impairs the ability of the metal to adjust to the strains imposed by static loads; moreover, the fatigue properties are degraded, lowering resistance to cyclic loads. When low-alloy steels are used for critical weldments in machinery, postweld heat treatments are employed to overcome most of these adverse effects. Preheat is necessary to permit these hardenable steels to overcome the effect of hydrogen, which can cause cracks in the HAZ during arc welding operations. Fusion Zone. In an arc weld, the fusion zone comprises the amount of base metal that has been melted in addition to the

filler metal that is fused by the arc. This is a cast structure distinctly different from the base material--which in wrought materials is usually a polycrystalline structure of more or less uniform size. The welding filler metals normally used for unalloyed carbon steels produce welds whose yield strength is generally higher in the as-deposited condition than that of the base metal. In most machinery applications, this higher strength is considered advantageous. Failures in static loading will therefore depend on the strength of the base metal, which the designer has chosen based on its known properties. Weld failures in statically loaded full-strength welds are thus unlikely, assuming that the welds meet specified workmanship standards. Low-alloy steel welds of matching weld-metal composition are considerably stronger than the base metal. In these applications, it is often preferable to choose a filler metal of lower alloy content (especially carbon) to match, or even undermatch, the strength of the base metal. As long as the yield strength of the weld metal is not less than that of the base metal, static loads will not cause significant plastic deformation in the welds. Because filler metals are traditionally classified on the basis of the ultimate tensile strength of the welds that they produce, it is possible to use filler metal that is one classification lower in ultimate strength than the base metal. Generally, the yield strength of the lower-strength weld metal will more nearly match that of the base metal. Table 20 lists matching filler metal specifications.

TABLE 20 MATCHING FILLER METAL SPECIFICATIONS

BASE METAL(B) CLASSES I AND II

WELDING PROCESS(A) SMAW SAW AWS A5.1 OR A5.5 AWS A5.17 F6X E60XX OR E70XX OR F7X-EXXXX

CLASS III

AWS A5.1 OR A5.5 E70XX(D) AWS A5.5 E80XX(D)

CLASS IV(E) CLASS V

AWS A5.5 E110XX(D)

AWS A5.17 F7X-EXXXX AWS A5.23 F8X-EXXX AWS A5.23 F11X-EXXX

GMAW AWS A5.18 E70S-X AWS A5.18 E70S-X GRADE E80S(F) GRADE E110S(F)

FCAW AWS A5.20 E60T-X OR E70T1 (EXCEPT EXXT-2 AND EXXT-3)(C) AWS A5.20 E70T-X (EXCEPT E70T-2 AND E70-T3)(C) GRADE E80T(F) GRADE E110T(F)

Note: The use of the same type of filler metal having the next higher mechanical properties as listed in AWS specification is permitted.

(A) A FILLER METAL OF A LOWER STRENGTH MAY BE USED WHERE THE DESIGN ENGINEER HAS STIPULATED LESS THAN 100% JOINT EFFICIENCY.

(B) IN JOINTS INVOLVING BASE METALS OF DIFFERENT YIELD POINTS OR STRENGTHS, FILLER METALS APPLICABLE TO THE LOWER-STRENGTH BASE METAL MAY BE USED. (C) IF TYPE 2 FILLER METALS ARE USED ON SUCCESSFULLY TESTED PROTOTYPE VEHICLES, THIS QUALIFIES THE USE OF TYPE 2 FILLER METALS FOR PRODUCTION. (D) LOW-HYDROGEN CLASSIFICATION. (E) WHEN WELDS ARE TO BE THERMALLY STRESS RELIEVED, THE DEPOSITED WELD METAL SHALL NOT EXCEED 0.05% V. (F) FOR WELD METAL HAVING A MINIMUM SPECIFIED YIELD STRENGTH GREATER THAN 415 MPA (60 KSI), THE USER SHALL DEMONSTRATE THAT EACH COMBINATION OF ELECTRODE AND SHIELDING PROPOSED FOR USE WILL PRODUCE LOW-ALLOY WELD METAL HAVING THE MECHANICAL PROPERTIES LISTED. THE MECHANICAL PROPERTIES SHALL BE DETERMINED FROM A MULTIPLE-PASS WELD MADE IN ACCORDANCE WITH THE TEST REQUIREMENTS OF THE LATEST EDITION OF AWS A5.18 ("SPECIFICATION FOR MILD STEEL ELECTRODES FOR GAS METAL ARC WELDING") OR A5.20 ("SPECIFICATION FOR MILD STEEL ELECTRODES FOR FLUX CORED ARC WELDING"), AS APPLICABLE. WHEN AN APPLICABLE AWS FILLER METAL SPECIFICATION IS ISSUED, IT WILL TAKE PRECEDENCE OVER ALL OTHERS, AND TESTING BY THE USER WILL NOT BE REQUIRED. THE MECHANICAL PROPERTY TESTS REQUIRED FOR GRADES E110S. E100T, AND E110T SHALL BE MADE USING ASTM A 514 BASE METAL. WHEN REQUESTED, THE ELECTRODE MANUFACTURER SHALL FURNISH CERTIFICATION THAT THE ELECTRODE WILL MEET THE REQUIREMENTS OF CLASSIFICATION OR GRADE. As-Welded Versus Postweld Heat-Treated Welds. Most welds used for machinery and equipment, including HSLA steels, are placed in service in the as-welded condition. Postweld heat treatment is usually specified to reduce residual stress and to attain dimensional stability when the weldment is to be machined after welding and the machining dimensional tolerances are small. Postweld heat treatment is sometimes specified to improve corrosion resistance in the welded area. The allowable stresses in specifications for components to be postweld heat treated are based on experience or on laboratory tests of the various types of weld joints.

High residual stresses remain in welds that have not undergone postweld heat treatment. In some cases, service loads added to the residual stresses in the weld area will cause sufficient distortion to present problems of dimensional misalignment. A stress-relief thermal treatment may be needed to overcome problems caused by residual stresses. Additional information is available in the article "Residual Stresses and Distortion" in this Volume. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Automobiles

Since its inception nearly a century ago, the automobile industry has relied heavily on joining technologies. Improvements in those technologies over the years have enabled improvements in automotive design--a constantly changing field driven by the ever-increasing demands of customers and governmental agencies. Customers want vehicles to meet criteria ranging from performance and function to appearance and styling; governmental agencies dictate that vehicles meet stringent safety and environmental standards. Producing automotive designs under such demands and constraints can often create difficulties in joining operations. Service Conditions. The automotive design engineer must consider numerous types of service conditions. Structurally,

a vehicle must be capable of withstanding on- and/or offroad service, must be crashworthy in order to meet Federal Motor Vehicle Safety Standards (FMVSS), and yet must possess an economy in terms of overall weight to meet Corporate Average Fuel Economy (CAFE) regulations. A vehicle must be able to withstand diverse environmental conditions,

ranging from the external environment (weather, electrolyte exposure, acid rain, etc.) to the internal environment (engine heat, fuels, lubricants, sealers, etc.), and must be watertight and airtight as well. Finally, all of this must be accomplished within the boundaries of cost-effective and manufacturable processes, including joining. Welding Processes. The automobile industry utilizes a wide array of welding processes, materials, coatings, and

designs. Given the many possible combinations of these factors, difficulties in the manufacturing process are inevitable. However, many problems can be avoided by exercising care in the initial design phase. By taking into account structural requirements (strength, crashworthiness, and fatigue resistance) as well as appearance requirements (corrosion resistance and finish), the design engineer can determine the most effective combination of materials and joining processes to fit the intended design and maximize performance. Joining processes commonly used in the automotive industry include, but are not limited to, the following: • • • • • • • • • • •

RESISTANCE SPOT WELDING GAS-METAL ARC WELDING GAS-TUNGSTEN ARC WELDING FLUX-CORED ARC WELDING RESISTANCE PROJECTION WELDING STUD ARC WELDING LASER BEAM WELDING RESISTANCE SEAM WELDING BRAZING SOLDERING PLASTIC WELDING

These processes are applied and controlled either by the automotive industry itself or by suppliers within the industry. In either case, each joining technique should be governed by a reliable set of standards to ensure that the technique is properly defined and applied. The sections that follow will describe how the more commonly used joining techniques listed above are affected by the various material and mechanical factors encountered in the automotive industry. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Resistance Spot Welding Resistance spot welding is the primary joining process used in automobile production. Its widespread use can be attributed to its many advantages. Specifically, the method is easy, cost effective, and, when properly applied, can produce structures with greater integrity than the sum of the individual welds. Therefore, RSW lends itself to the extremely high production volumes common in the automotive industry. Despite the obvious advantages of the spot welding process, adequate results cannot be achieved without special consideration of a number of factors that can ultimately affect the welded structure, including raw material coatings, processing fluids, raw material chemistry, joint design, joint sealers, and joint adhesives. Steel Composition. Raw material chemistry is of primary importance in the design of a spot-welded structure. Most

automotive applications use SAE 1008-1010 steel. This material has a low carbon equivalent, which makes it readily formable and weldable and relatively inexpensive and abundantly produced. However, low-carbon material is not used exclusively. Automobile manufacturers must reduce weight while retaining structural integrity. This has resulted in increasing use of high-strength steels, such as bake-hardenable, dual-phase, precipitation-hardenable, and martensitic steels. These materials are inherently more difficult to weld because of the residual elements added for solid-solution strengthening, but weldable grades are available. As a rule, welding of these materials requires increased heat input compared with low-carbon steel material.

Steel Thickness. In the design of a spot-welded joint, a critical factor is the combination of materials to be joined, along

with their thicknesses. As the number of sheets to be welded increases, the difficulty of the welding operation also increases. Similarly, welding becomes more difficult as the ratio of material thicknesses between sheets becomes overly large. Flange size is another important factor. Flanges must be large enough to accommodate weld equipment constraints and required weld sizes. Too often, designers fail to consider these important yet simple manufacturing requirements. Coated Steels. Due to consumer demand, the automobile industry has been striving in recent years to produce vehicles

with greater corrosion resistance. For the most part, this demand has been met by the use of steels coated by galvanizing, aluminizing, prepainting, or laminating. Such coatings greatly enhance the corrosion resistance of the base material, but reduce its weldability. In general, spot welding of coated materials generally requires an increase in weld current and hold time. Greater electrode wear is also exhibited, which can ultimately affect the resultant welds. Joint Sealers. Passenger comfort is always an overriding concern in the design of a vehicle; therefore, much attention is

paid to the reduction of wind noise, fume intrusion, and water leakage. Body sealers, particularly weld-through sealers, are used to control these factors. Welding problems associated with the use of the sealers include reduced conductivity at the faying surface and increased fuming as the sealer is burned away at the weld. Solutions to these problems include increasing heat input to ensure proper welding and providing for adequate ventilation to remove fumes generated (although minimal) by the vaporized sealer. Joint adhesives are used to enhance joint integrity. They are commonly used in conjunction with the spot welding

process to improve joint shear strength. Similar to joint sealers, adhesives can cause problems in welding, which can be overcome by the methods described above. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Gas Metal Arc Welding The GMAW process is widely used in the automotive industry to weld frames, brackets, body panels, and many other structural components. This method is easily applied (automated and manually) and lends itself well to high productivity at a relatively low cost; in addition, it allows most metals to be joined to themselves. Disadvantages include the fact that GMAW is sensitive to material chemistry, surface coatings, and joint design. When these factors are not considered in the initial design, weld quality problems, such as cracks, porosity, lack of fusion, and burnthrough, often result, which may affect the structural integrity of an assembly. Steel Properties. In the automotive industry, a wide variety of steel chemistries are used to meet strength and durability requirements, ranging from low-strength plain carbon steels to high-strength alloys. To avoid GMAW problems associated with chemistry, the designer must take into account the compatibility of the materials to be joined as well as their individual chemistries. Both plain carbon and alloy steels with carbon equivalents of approximately 0.30 are generally considered to be weldable. Once the material chemistries have been established, the appropriate solid filler material must be chosen to produce a weld with characteristics similar to those of the base materials being joined. By taking these initial design steps, problems such as underbead cracking and weld-metal embrittlement can be avoided. Surface Condition. The GMAW process is sensitive to surface impurities on the base material, such as corrosion, dirt,

and oil. Surface coatings can also adversely affect the resultant weld. When improperly prepared surfaces are welded by this method, quality problems (typically porosity and cracking) are often observed. The solution involves simply removing the surface contaminants. Surface coatings are not dealt with as easily, for either the coating must be selectively removed from the weld joint or the welding parameters must be adjusted to reduce the effects of the off-gassing material. This is sometimes done through a trial-and-error process, but can be refined to produce sound welds. In addition, if the coatings are removed in the weld area, provisions must be made to reapply corrosion protection where necessary. Joint Preparation. Joint designs for GMAW are quite important in the automotive industry. Joints must be designed in

such a way as to minimize the amount of joint preparation necessary, yet produce a weld that will serve the intended purpose. Joint preparation is a time-consuming process, which is not desirable for high-volume vehicle production.

Therefore, naturally occurring joints that result from strategic part placement are preferred to prepared joints, such as groove welds, which often require extra processing steps. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Gas Tungsten Arc Welding and Flux-Cored Arc Welding The GTAW process is well suited to handle thin-section welding with or without filler materials. The FCAW process can handle the same types of welding conditions as the GMAW process, despite the filler material differences (flux-filled wires versus solid wires, respectively). In terms of automotive applications, both processes involve considerations similar to those listed for the GMAW process. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Fasteners: Resistance Projection Welding and Stud Welding The intricate assembly of an automobile requires that parts be securely fastened, yet the ease of manufacturing must not be compromised. In order to accomplish this task, projection-welded screws, nuts, and studs are often used. These fasteners must possess welds that can withstand the rigors of a production environment as well as endure the torquing operations in assembly. Types of service conditions that can affect the welding and subsequent use of these parts include: • • • •

BASE METAL/FASTENER CHEMISTRY BASE METAL/FASTENER STRENGTH CHARACTERISTICS BASE METAL/FASTENER COATINGS BASE METAL CLEANLINESS PRIOR TO WELDING

Materials Properties. Selection of the correct base metal and fastener chemistries is extremely important to ensure

compatibility at the weld interface and HAZs. In order to avoid exotic weld schedules, materials with carbon equivalents of approximately 0.30 are ordinarily used, as they are readily weldable. It should be noted that these types of welds result in a similar type of weld metal/HAZ transition as other joining processes; this fact must be recognized in the design of such weldments. Base-metal strength and fastener strength characteristics can affect the ultimate mechanical joint integrity. Designs must reflect conditions where the base-metal strength is compatible with that of the fastener; otherwise, premature failures or insufficient mechanical joints may result. The different strength characteristics may be a result of chemistry, heat treatment, cold working, or simply section thickness. Surface Condition. Projection-welded fasteners may also be affected by coatings present on the base material or the

fastener. As described previously, the automotive industry utilizes different corrosion-resistant coatings on many areas of vehicles. Such coatings can affect projection or stud welds in several ways. If the base material is zinc coated or the fastener is coated with a corrosion-resistant plating, the arc characteristics of the welding process are much different than for bare materials. In either case, provisions must be made in welding schedules. Lastly, use of an antispatter/antifriction coating, such as Teflon, can also affect the welding process. Such coatings are being used increasingly to prevent production problems associated with weld spatter and expulsions generated by other processes. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Laser Beam Welding Laser beam welding is an emerging technology that has shown promise for use in the automotive industry. Basically, the technique utilizes a high-power laser beam capable of penetrating deep into the weld joint. Laser welding offers high speed, precision, and flexibility to automotive applications. The resultant welds exhibit very small HAZs and are easily formed to final part configuration without weld degradation. The diverse capabilities of the laser welding technique enable the user to produce parts with properties unattainable with common methods, such as resistance spot welding or GMAW. Laser welding is most often used to make lap and butt joints. It allows blanks to be fabricated through the welding of different-gage materials as well as materials with dissimilar coatings. After the welding operation, the parts can be formed to their final shape. This allows heavier-gage materials to be used in place of reinforced areas, thereby eliminating parts. Similarly, coated parts can be strategically placed in corrosion-prone areas, which eliminates the need to coat the entire assembly. Despite the obvious advantages of laser welding, application of the technique requires care, especially with the types of materials used in automotive applications. Coated steels complicate the laser welding process, because the off-gassing of the lower-boiling-point material interferes with the maintenance of the weld keyhole. Once this keyhole is broken, the weld must be restarted, resulting in inferior quality. Joint edges and part alignment/fit-up are also common production problems when laser welding. They affect the continuity of the welding process by disrupting the laser beam or by deflecting the beam from the joint. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Other Joining Processes Many other joining techniques are commonly used in the automotive industry, but to a lesser extent than those discussed above. Some of these processes include resistance seam welding, soldering, brazing, and plastic welding, each of which requires due care in its application to ensure the production of sound joints. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Railroad Equipment

The transporting of freight and passengers by railroads in the United States involves over 200,000 miles of track, 1,200,000 freight cars, 19,000 locomotives, and thousands of passenger vehicles (Ref 19). Welding is the major process used in the fabrication of railroad facilities and structures, as well as in their repair and maintenance. Practically all known welding processes are being used or have been applied to components used by the industry. This section will deal with freight cars, locomotives, and track.


19. RAILROAD FACTS, ASSOCIATION OF AMERICAN RAILROADS, 1990


Freight Cars Freight car requirements vary considerably, depending on the type of load being carried. A variety of products are transported by rail, including various liquid and solid commodities. Some commodities must be kept at low temperatures and some at high pressures, and some are classified as hazardous. Public safety is always a key issue. Many freight cars are in "interchange service," which means that they may be handled by any railroad. It also means that repair work may be done by a variety of shops throughout North America. Ease of repair welding is therefore a critical consideration in material selection. It is not unusual to expect a freight car to last 40 years. Cars for interchange are designed for required service conditions, which for high-utilization cars may be up to 3 million miles of service. Freight cars must be able to sustain 4.5 MN (106 lbf) of squeeze and single-end impacts of 5.6 MN (1.25 × 106 lbf). Designs must include buff and draft forces of up to 3.5 MN (630,000 lbf) without exceeding the yield strength of the materials. Freight cars are subject to a variety of cyclic stresses due to train action, track irregularities, and track hunting. Most failures are caused by fatigue crack propagation, with a lesser number caused by a single overspeed impact. Types of freight cars in common use include tank, flat, gondola, box, hopper, and covered hopper cars. Of these, only

tank and covered hopper cars have seen a net increase in number over the last 21 years (Ref 20). The largest increase has been in the use of covered hopper cars, which handle a variety of dry commodities that must be kept clean. Tank cars consist of a tank, which is basically a pressure vessel, and an underframe, which provides connection to other cars and to the "trucks." Underframes may be continuous or may consist of attachments to the ends of the tank, with the tank being a structural member. Tank cars and covered hopper cars are often lined with epoxy, rubber, or other protective materials to prevent car body corrosion and/or contamination of the commodity being carried. Specifications. Freight car construction is regulated by the Association of American Railroads (AAR). Tank car welding

is covered by the AAR "Tank Car Specification." For other types of freight car welding, the AAR references the AWS "Railroad Welding Specification" (D15.1). Design Considerations and Materials of Construction. With the type of welded construction employed on freight

cars, the use of higher-strength materials does not improve fatigue resistance. Where fatigue is a significant factor, structural designs incorporating welding are based on carbon steel and HSLA steels with moderate strengths (usually 345 MPa, or 50 ksi, yield strength or less). For applications requiring high static strength and moderate fatigue resistance, quenched and tempered steels with yield strengths of 690 MPa (100 ksi) and high-strength controlled rolled steels with yield strengths of 550 MPa (80 ksi) have been used. The most common steels used in the production of modern tank cars are ASTM A 516 grade 70 and AAR TC128 grade B (similar to ASTM A 612 grade B). For all pressure tank cars and for some nonpressure cars, these steels are used in the normalized condition. For general-purpose nonpressure cars, hot-rolled steel is used. Certain components are exposed to severe wear conditions, including "truck" components as well as couplers and related wear plates. Wear materials used include plain high-carbon steel, quenched and tempered alloy steel, and austenitic manganese steel, which are attached by welding or by mechanical fasteners. Some commodities warrant the use of stainless steel tanks or covered hopper cars, which are generally fabricated of AISI type 304 or 316 stainless steel. However, this constitutes a small percentage of the freight car business. The empty car weight, or tare weight, is an important factor in determining the hauling capacity of a car. The tare weight is deducted from the total allowed weight (i.e., gross rail load) to determine the weight of commodity that can be loaded (i.e., load capacity). The energy used to haul the tare weight around is wasted fuel. For this reason, the use of aluminum

car bodies has recently become more popular, especially in coal-hauling service. The corrosion resistance of aluminum alloys is an added benefit. In some designs, the aluminum is mechanically fastened to the welded steel underframe, and little or no welding of the aluminum is required. Freight cars generally consist of six components: tracks, brakes, underframes, draft components, safety appliances, and bodies. Figure 13 shows a side view of a general-purpose boxcar. A high percentage of the welding operation is performed in the underframes and the car bodies. Underframe components typically are from 6.4 to 25 mm ( thick, but may be thicker for heavy-duty cars. Bodies vary from 3.2 to 7.8 mm ( range from 11 mm (

1 to 1 in.) 4

1 5 to in.) in thickness. Tank car tanks 8 16

7 in.) to slightly more than 25 mm (1 in.) in thickness. 16

FIG. 13 SIDE VIEW OF A GENERAL-PURPOSE BOXCAR. 1, END SHEET (TOP); 2, END LINING (TOP); 3, END SHEET (BOTTOM); 4, END LINING (BOTTOM); 5, END SILL; 6, STRIKER FACE RIB; 7, STRIKER FACE; 8, KEY SLOT LINER AND FRONT DRAFT LUG; 9, COUPLER CARRIER; 10, DRAFT GEAR; 11, FRONT DRAFT GEAR STOP; 12, REAR DRAFT LUG RIBS; 13, COMBINATION CENTER FILLER/CENTER PLATE; 14, TRUCK FRAME; 15, NAILABLE STEEL FLOORING; 16, WHEEL; 17, FLOOR BEAM; 18, LADING STRAP ANCHOR; 19, CENTER SILL; 20, CROSSBEARER; 21, FLOOR STRINGER; 22, DOOR FRONT STOP; 23, DOOR POST; 24, DOOR LOCK BRACKET; 25, DOOR LOCK; 26, DOOR LOCK REINFORCING GUSSET; 27, DOOR ROLLERS; 28, DOOR POST AND SIDE SILL GUSSET; 29, STARTER AND CLOSURE LEVER; 30, ROUTING BOARD; 31, PLACARD BOARD; 32, DOOR HANDLE; 33, FORKLIFT PUSHER BLOCK; 34, DOOR CATCH; 35, SPARK STRIP; 36, SIDE SILL REINFORCEMENT; 37, DOOR TRACK BRACKET; 38, SIDE POST; 39, DOOR TRACK (BOTTOM); 40, DOOR CATCH ASSEMBLY; 41, SIDE SILL; 42, AXLE; 43, TRUCK SPRINGS; 44, TRUCK BOLSTER; 45, JACKING PAD; 46, SILL STEP; 47, UNCOUPLING ROD; 48, KEY SLOT; 49, CROSSOVER STEP; 50, HAND BRAKE WHEEL; 51, LADDER GRAB; 52, LADDER STILE; 53, GRAB; 54, CORNER POST; 55, SIDE PLATE; 56, ROOF SHEET (END); 57, ROOF SHEET (INTERMEDIATE); 58, SEAM CAP; 59, SIDE SHEET (PLAIN); 60, DOOR SAFETY HANGER; 61, SLIDING DOOR; 62, LIFTING LUG; 63, DOOR TOP RETAINER; 64, ANTIPILFERAGE LOCK. COPYRIGHT SIMMONSBOARDMAN PUBLISHING COMPANY/THE RAILWAY EDUCATIONAL BUREAU. REPRODUCED WITH PERMISSION

Welding Processes. The most common welding processes used in the manufacture of freight cars are FCAW, GMAW, and SAW. SMAW has been largely displaced by semiautomatic, mechanized, or automatic welding procedures. Seams in tank car tanks are produced primarily by SAW. Except for carbon and low-alloy steel tank car tanks, which are stress relieved, most freight car components are used in the as-welded condition.


20. FREIGHT CARS 1992: THE QUALITY PROCESS AT WORK, PROGRESS. RAILROAD., APRIL 1992 Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Locomotives Modern locomotives are diesel-electric, ranging from 0.75 to 3.3 MW (1000 to 4500 hp) in power and weighing as much as 181,000 kg (400,000 lb). The life of a locomotive is expected to be 20 years and with rebuild can be extended to as long as 40 years. Locomotives are generally owned by a railroad and remain within its control. Maintenance is normally done by the owners, but major overhauls and rebuilds are performed by remanufacturing specialists. Design Considerations and Materials of Construction. The modern diesel-electric locomotive consists of welded

components fabricated using the full range of welding processes. The basic components are the underframe, cab structures, engine, generator, tracks, traction motors, electrical cabinets, fuel tank, air compressor, and brake equipment. Components are fabricated from rolled plate, sheet metal, forgings, castings, and various shapes. Locomotives must be able to withstand forces similar to those experienced by freight cars, including 4.5 MN (106 lbf) of squeeze. Buff and draft forces are transmitted through a self-supporting rigid underframe consisting of a fabricated draft gear pocket, "T" or "I" sills, and a bottom plate. Materials used for fabrication of the main components of a locomotive are typically selected on the basis of good weldability in order to facilitate original construction and subsequent maintenance and rebuild. The underframe, fuel tank, traction motor frame, generator frame, and other structural steel components are fabricated from low-carbon steel containing less than 0.26% C, with and without added alloying. Typical tensile and yield strengths are 415 and 275 MPa (60 and 40 ksi), respectively. Weld filler metals used for these components include E70xx, ER70S-X, and E70T-X-type electrodes. Although higher-strength base materials do not improve high-cycle fatigue life, quenched and tempered steels with a tensile strength of 690 MPa (100 ksi) and normalized carbon-manganese steels with a tensile strength of 550 MPa (80 ksi) may be used when high static strength or weight reduction is an issue. Diesel locomotive track frames are made of cast steel. Except for the application of wear plates, welding on cast steel truck frames is accomplished using E70-type electrodes. Wear plates made of medium-carbon, high-carbon, or austenitic manganese steels are welded using austenitic stainless steel electrodes. With the exception of certain critical subassemblies (bolsters, center plates, and traction motor frames), which are stress relieved, most diesel locomotive structural fabrications are used in the as-welded condition. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Track Rails are designed for strength and wear resistance. Locomotives and freight cars weighing up to 181,000 kg (400,000 lb) and riding on steel wheels create high bearing stresses, requiring very-high-strength rail material. Bearing stresses may be up to 2750 MPa (400 ksi). Rail life averages 70 years. Replacement of rail is extremely expensive. Maintenance of high wear areas, such as switch points and rail ends, requires that repairs be made economically in the field. The desire for long, continuous rails necessitates the use of materials that allow butt welds to be made in the rail that have properties matching the unwelded rail.

Design Considerations and Materials of Construction. There are two categories of rail materials. The largest

category by far is T-rail, typically of near-eutectoid composition (~0.8% C), which makes up the thousands of miles of railroad track. The much smaller category is austenitic manganese steel, which is used extensively in cast frogs. The welding practices for each type of material are very different. This section will address the metallurgical reasons for these differences. Welding is applied to T-rail for joining and repair purposes, and is applied to austenitic manganese steel solely for repair purposes. The controlling feature for welding of near-eutectoid T-rail is hardenability--that is, its tendency to form transformation products such as martensite and bainite, rather than the preferred pearlite. Untempered martensite is unacceptable because of its possible embrittling effect on rail steel; near-eutectoid composition bainite and tempered martensite are avoided because of their poor wear resistance compared with pearlite of the same hardness. Historically, rail steels have had carbon contents near 0.7 to 0.75%, with manganese contents of 0.85 to 0.95% and silicon levels of up to 0.5% maximum. The as-hot-rolled hardness generally has been near 240 to 270 HB. More modern conventional carbon rail steel has a slightly higher maximum manganese content (1.25%) and allows up to 0.25% Cr to be added to ensure as-hot-rolled hardness near 300 HB. Metallurgical Characterization. The time-temperature-transformation (TTT) characteristics of steel control the type of

transformation product that results from the welding process. A typical continuous cooling curve diagram for the older grades of conventional carbon rail steel is illustrated at the left in Fig. 14. In an effort to ensure that the transformation product resulting from welding is pearlitic, the maximum cooling rate through the transformation region generally must not exceed approximately 4 °C/s (7 °F/s). For rates up to this maximum level, transformation will take place entirely between the pearlite start and finish boundaries. However, should the cooling rates exceed approximately 4 °C/s (7 °F/s), it is possible that the cooling path will miss the nose of the pearlite finish boundary such that untransformed austenite will be available for transformation to bainite or untempered martensite at temperatures below 540 °C (1000 °F).

FIG. 14 TTT CHARACTERISTICS OF RAIL STEELS

Alloying additions are sometimes made to near-eutectoid rail steels in an effort to produce pearlite having a smaller spacing between the Fe3C platelets during cooling following hot rolling; this yields superior hardness and greater wear resistance. The alloying additions most commonly used are the carbide formers: chromium, molybdenum, and vanadium. If added in sufficient amounts (1% Cr; 0.6 to 0.8% Cr and 0.20% Mo; or 1% Cr and 0.08% V), these types of alloy additions tend to increase hardenability significantly and in so doing shift the diagram to longer times. In fact, a specific bainite nose can be obtained (see Fig. 14).

Welding Thermal Cycles. One of the limitations of the alloy approach to strengthening is that a welding practice

resulting in a cooling rate near 2 °C/s (3.6 °F/s) (which would be perfectly satisfactory for conventional carbon steel rail) would now lead to the formation of undesirable bainite or untempered martensite. Welding practices must be modified to slow the cooling rate to less than 2 °C/s (3.6 °F/s). By and large, most rail welding is done to join rails into strings up to 440 m (1440 ft) in length. The most common practice for that purpose is flash butt welding utilizing either ac or dc current. This practice generally involves the use of a number of preheating cycles to heat the rail ends and control the subsequent cooling rate. When the rail ends have been sufficiently preheated, a flashing phase is initiated that removes oxidized metal and establishes a metal ion plasma in which forge welding of the rail ends can be accomplished. The forging operation ideally expels all liquid metal from the bond interface, such that the weld is in reality a solid-state bond. Postweld Heating and Cooling Cycles. For significantly alloyed rail steels, adjustment of the preheat practice may not

be sufficient to limit the final cooling rate. To this end, postheat cycles are applied after the forging cycle. The added expense of postheating has caused railroads to prefer the use of heat-treated carbon steel rails--both those achieving full section hardening and those in which only the head is hardened. These rails generally require no modification of the electric flash butt welding practice. However, a consequence of welding most heat-treated rails (except those classified as slack quenched) is the loss of the heat-treat strengthening in the weldment. This loss can lead to severe batter of the weldment under heavy wheel loads. Thus, some railroads have introduced a postweld air quench to refine the pearlite in weldments. An alternative approach to air quenching is the use of head-hardened rails having their composition adjusted slightly to achieve slack quenching during the normal cooling after flash butt welding. Typically, such modified compositions will include increased silicon (up to 0.8%) and chromium (up to 0.5%) levels. After normal cooling, the hardness of the weldment will have returned to about the same level (~350 HB) as that of the head-hardened region of the rail. Although this approach works very well, the cost of the rail may be notably higher than that of conventional head-hardened rail. Neither air quenching nor slack quenching prevents the softening that occurs in the outer boundaries of the HAZs. Other Welding Processes. Rails are also joined together in track by thermite welding--a metal casting process based on

the reaction between aluminum powder and iron oxide (see the article "Thermite Welding" in this Volume). The resulting cast structure has low ductility and toughness compared with the hot-worked structure of a flash butt weld. The cooling rates typical of thermite welding are low in comparison with those of the flash butt process, so there generally is no need to alter standard welding practices for use with alloy rail. Welding also is applied to near-eutectoid rail steels to repair battered and chipped rail ends, wheel burns, and worn and chipped switch points and frogs. Both arc welding and oxyfuel gas welding processes are possible, although arc welding is preferred. Either coated electrodes or flux-cored wire can be used. Electrodes may contain high nickel and manganese contents, and deposits are generally austenitic. Both longitudinal and transverse lay-down patterns have been employed. Austenitic Manganese Steels. The nominal chemical composition range for austenitic manganese steel is 10 to 14%

Mn and 1.0 to 1.4% C. The presence of such high levels of manganese stabilizes the high-temperature fcc austenite phase so that, with rapid cooling, the austenite is retained to room temperatures. The austenite has a relatively low yield strength that is not strongly influenced by carbon or manganese levels. However, its ultimate tensile strength, which is influenced by carbon or manganese levels, can be quite high in comparison with yield strength. This wide separation imbues the steel with a much higher work-hardening capacity than pearlitic rail steels. It is this high work-hardening capacity that makes manganese steels effective in resisting impact loads such as those occurring at railroad frogs. Although the high manganese content can delay the transformation of the high-temperature austenite phase to the room-temperature equilibrium phases (ferrite and carbides), it cannot completely suppress the transformation. If temperatures are elevated to 315 °C (600 °F) for very long periods or to temperatures 480 to 540 °C (900 to 1000 °F) for much shorter periods (a few minutes), the steel can become embrittled. The mechanism of embrittlement is the precipitation of carbides, usually along austenite grain boundaries. For this reason, repair welding must cease when the casting temperature reaches 315 to 370 °C (600 to 700 °F), and the parts must be allowed to cool to a more suitable range.


Sheet Metals This section deals with a wide variety of sheet metal fabrications, excluding those discussed in previous sections of this article. Ductwork of many types, small and large appliances, containers, pharmaceutical and food processing vessels, and industrial and domestic kitchen equipment are examples of sheet metal applications. Service conditions are the principal consideration in material selection for such applications. The material attributes

listed in Table 1 are a useful starting point. If an application requires cold forming or deep drawing, the ductility of the material will be of primary importance. Corrosion properties are paramount in applications where the weldment will not be protected by some type of coating following fabrication. Mechanical properties figure prominently in designs for structural applications. Density becomes important in many applications for products where shipping weight is a factor. Fracture toughness, conductivity, and magnetic properties are rarely significant in sheet metal fabrications. Weldability is extremely important in the material selection process, particularly when an application involves very thin

sheet materials. The choice of welding process will be determined by the ability to weld thin materials without burnthrough. When dealing with sheet metals, the ability to make sound welds without burnthrough is often synonymous with weldability. The weldability of sheet metals is vitally affected by the welding process selected, the manner of welding (manual, semiautomatic, or automatic), and the number of parts to be fabricated in a given manufacturing operation. This section will deal with material selection criteria as determined by the joining process to be employed. Welding process selection is greatly influenced by the need to control distortion. Thinner material is inherently more difficult to weld because of the thermal cycles encountered in welded joints. Distortion is minimized by rigid fixturing and by use of high-energy rapid-travel-rate welding processes. Brazing or soldering allows lower peak temperatures to be used for joining, which is advantageous in distortion control. Materials with low heat conductivity, such as stainless steels, aggravate distortion problems. Warpage of materials during joining is one of the principal problems encountered in the fabrication of sheet metal by any of the joining processes that involve fusion of metal. Material Requirements for Service Conditions ASM Committee on Material Requirements for Service Conditions*

Joining Processes Resistance Welding. Spot, seam, and projection welding are commonly used in sheet metal fabrications. The most

commonly used material joined by these processes is low-carbon steel, although low-alloy steels, stainless steels, nickel, and copper-base alloys are also readily joined. When an unalloyed carbon steel is suitable for an application, the carbon content is generally limited to 0.20% maximum. For products that are to be enameled after welding, the carbon content is usually kept below 0.04% to avoid blistering of the enamel. The designer can then take advantage of the enamel thickness in assessing the strength properties of the part, which will offset the relatively low strength of the very-low-carbon steel. Spot welding and projection welding are advantageous for control of distortion, but are unsatisfactory for containment of fluids or for elimination of crevices. Seam welding can be suitable for fluid containment, but distortion is a greater problem and crevices continue to be present. As described in the section titled "Automobiles" in this article, joint sealants and adhesives can be used with the resistance welding processes, offering the possibility of fluid containment and crevice elimination. Arc Welding. Developments over the past 20 years have greatly increased the application of arc welding in the

fabrication of sheet metals. All of the arc welding processes are employed, although some are more suitable than others.

The simplest arc welding process is SMAW. For nonrepetitive parts, it is the preferred process and is suitable for thicknesses of most steels down to about 2 mm (0.080 in.). Covered electrodes down to 1.6 mm (

1 in.) in diameter are 16

available in most types of low-carbon, low-alloy, and stainless steels. Burnthrough is a problem with the thinner sheet metals, more so with stainless steel than with carbon steels because of its low thermal conductivity. The SMAW process is commonly used in job shops and for repairs. Fixturing for prevention of distortion and warpage is needed in cases where large, flat surfaces must be kept flat. Manual welding with the gas-shielded arc welding processes (GMAW, FCAW, and GTAW) offers improved control of the arc energy, especially when using more sophisticated power sources with pulsed current controls. However, travel rate in manual welding is generally slow, so the distortion problem is similar to that encountered with SMAW. Automation of the process offers increased welding speed, use of fixtures with heat sinks for more rapid cooling, and higher productivity and quality. High-energy-beam welding processes, such as electron and laser beam, offer even higher welding speed and lower distortion. Materials sensitive to atmospheric effects during welding, such as titanium and zirconium, are especially suited to electron beam welding in vacuum chambers or with an inert gas trailing shield. Laser welding is being increasingly used for unalloyed carbon steel sheet fabrication, because the process requires no protection from the atmosphere during welding and because relatively low-power laser equipment of nominal cost is adaptable for thin metal sections. Fully automatic welding systems are generally required for high-energy-beam welding processes. Brazing and soldering are suitable joining methods for many sheet metal materials. The metal surfaces must be

compatible with the liquid-phase joining material. For successful joining, both brazing and soldering depend on capillary action between the mating surfaces to cause flow of the liquid joining material. Oxides on the metal surface must be mechanically or chemically removed in many instances for the liquid metal to flow freely. Fluxes are frequently used to assist this flow process. These joining processes are easily adaptable to manual low-volume fabrication using oxyfuel torches for the heat source. Job shops customarily use these methods for fabrication or repair. The lower joining temperature is a distinct advantage in that fusion of the base metals does not take place, minimizing distortion. Soldering is commonly used to join galvanized steel and tin plate. The surface zinc or tin becomes alloyed with the solder during the fusion process and assists in establishing an adherent bond with the underlying steel. The coated surface adjacent to the joint is unaffected and is thus resistant to the service environment. Brazing requires much higher temperatures and offers many more types of filler metals than soldering. A filler metal must be chosen not only for its ability to flow at the brazing temperature, but also for its compatibility with the service environment. The filler metal is distinctly different from the base metal and may set up electrochemical cells, causing localized corrosion adjacent to the joint. Brazing and soldering are superior to resistance welding for producing crevice-free joints, but are usually less satisfactory than the arc welding processes for joints that require mechanical and corrosion-resistance properties comparable to those of the base metals. Nonfusion Joining Processes. Bonding of metal surfaces can be accomplished in many metal systems without fusion.

Diffusion bonding, ultrasonic welding, and similar techniques can be used to join very thin metals. These processes are particularly appropriate for the precious metals--for example, in making contacts to foils in electronic circuitry. They are economically unsuitable for continuous seams in liquid or gas containers.


Materials and Properties Steels. Low-carbon steels are characterized by their considerable ductility at ambient temperatures, which becomes even

higher at elevated temperatures. During the welding thermal cycle, plastic deformation occurs in the HAZ. Unless restrained by clamping devices or by the inherent shape of the piece, unwanted distortion in the weld area will result. The thinner the steel, the greater the distortion problem. Fusion welds have greater strength and lower ductility than the base metal. If a sheet metal is to be cold formed after joining, this strength mismatch may need to be addressed in the design. Although the ductility of the weld is less than that of the base metal, it is usually adequate for most cold-forming operations. Distortion that forms during the welding operation can sometimes be ignored if the component is subsequently cold formed. In corrosive or high-temperature environments, carbon steel weldments will perform in a manner generally comparable to that of the unaffected base metal. The cast structure of the weld is often less resistant to some environments than the wrought structure of the unaffected steel. Also, the welding process usually results in fine oxide inclusions in the weld, which are often less resistant to corrosion. The higher strength of low-alloy steels does not offer the same advantages to sheet metal welding applications that it does to structural steel applications. Thinner sections might be possible by using the higher-strength steels, but distortion and weldability considerations usually make such choices uneconomical unless weight savings are critical. Table 21 lists the principal ASTM specifications for weldable carbon and low-alloy sheet steels.

TABLE 21 PRINCIPAL ASTM SPECIFICATIONS FOR WELDABLE CARBON AND LOW-ALLOY SHEET STEELS

TYPE OF STEEL

PURPOSE

ASTM NO. COLD HOT ROLLED ROLLED LOW CARBON COMMERCIAL QUALITY A 366 A 569 A 794 A 659 STRUCTURAL A 611 A 570 ENAMELING A 424 A 424 DEEP DRAWING A 619 A 621 A 620 A 622 CARBON-MANGANESE ATMOSPHERIC CORROSION A 606 A 606 PRESSURE VESSELS ... A 414 LOW ALLOY HIGH STRENGTH ... A 715 NIOBIUM OR VANADIUM ADDED A 607 A 607 Aluminum and aluminum alloy specifications for sheet metal are included in ASTM B 209. The low density of aluminum when compared with most metals makes aluminum sheet metal desirable for fabrications where weight savings are important. It is also useful for its adherent oxide, which makes it resistant to many environments at ambient temperatures.

The refractory oxide on the surface of aluminum becomes increasingly important as the thickness of the sheet is reduced. Gas-shielded arc welding processes, especially those using ac power, are essential to break down this adherent refractory aluminum oxide in order to make successful welds. Manual welding with covered electrodes is employed occasionally; oxyfuel welding is not suitable. Brazing and soldering are possible only with chemically active fluxes that dissolve the oxide layer.

The high thermal conductivity of aluminum can cause distortion. Gas-shielded arc welding processes overcome this problem by concentrating the arc energy sufficiently to cause fusion to occur. Electron beam welding is also suitable. Aluminum alloys offer higher strength, which is especially important in applications where weight saving is a factor. Furthermore, their thermal conductivity is lower than that of pure aluminum, making fusion welding easier. Distortion control is usually provided by rigid fixturing and backups. The metallic backup of the fixtures offers the advantage of heat sinks and backside weld contour control. Adding an inert backing gas often prevents oxidation of the molten weld pool. Stainless Steel. The corrosion resistance of stainless steel is provided by chromium in an amount of at least 12%. When

the content of nickel is greater than 8%, the austenitic form of iron becomes stable at ambient temperatures. The most frequently employed stainless steels for sheet metal fabrication contain 18% Cr and 8% Ni and are known as austenitic stainless steels. ASTM specifications normally used in stainless steel sheet metal applications are listed in Table 22.

TABLE 22 PRINCIPAL ASTM SPECIFICATIONS FOR WELDABLE STAINLESS STEEL SHEET

TYPE PURPOSE ASTM NO. OF STEEL AUSTENITIC CHROMIUM-NICKEL GENERAL REQUIREMENTS A 480 GENERAL PURPOSES A 167 A 240 ARCHITECTURAL A 666 ROLL CLAD A 264 DUPLEX CHROMIUM-NICKEL GENERAL PURPOSES A 240 FERRITIC CHROMIUM GENERAL REQUIREMENTS A 480 GENERAL PURPOSES A 240 A 276 ROLL CLAD A 263 The austenitic stainless steels are extremely ductile and easily welded. The weld strengths, both yield and ultimate tensile, are generally higher than those of the base metal, but ductility remains high, permitting subsequent cold forming or straightening if necessary. The thermal conductivity of stainless steels is about 30% less than that of carbon steels. The thermal cycles during welding cause plastic deformation to occur in the widely heated zones adjacent to the weld. Moreover, the coefficient of thermal expansion for stainless steels is about 30% greater than that of carbon steels. These two properties account for the even greater distortion problems encountered in the austenitic stainless steels compared with carbon steels. Therefore, rigid fixturing and heat sinks are essential when welding these alloys. Stainless steels are principally used for their resistance to aqueous corrosion and for their strength and oxidation resistance at elevated temperatures. With regard to certain corrosive media, it is important to recognize that the thermal cycles during welding cause carbide precipitation to occur in the HAZ, with severe loss in corrosion resistance. This can be overcome by the use of low-carbon types (0.03% maximum) or those stabilized with titanium or niobium. The stabilized stainless steels are also stronger at elevated temperatures. The fluidity of the very-low-carbon stainless steels often affects the attainment of the desired weld profiles. This is particularly troublesome in autogenous GTAW when sulfur is unusually low (0.15%) will produce greater hardness and, therefore, an increased susceptibility to cracking. Thus, type 440A, B, and C stainless steels are not usually considered for applications that require welding, and filler metals are not readily available in type 440 compositions. In addition to the problems that result from localized stresses associated with the volume change upon martensitic transformation, the risk of cracking will increase when hydrogen from various sources is present in the weld metal. Therefore, the use of low-hydrogen consumables is mandatory. Hydrogen-induced cold cracking in martensitic stainless steel weldments is also discussed in more detail later in this article. Selection of Wrought Martensitic Stainless Steels J.R. Davis, Davis & Associates

Weld Microstructure (Ref 2, 3, 4) Detailed studies have been made on the transformation, tempering behavior, and resulting microstructures of martensitic stainless steel weldments. These studies were spurred by an interest in using 12Cr-1Mo-0.3V (HT9) martensitic stainless steel as a candidate material for the first wall/blanket of fusion-containment devices (Tokamaks), because of its improved resistance to swelling, low coefficient of thermal expansion (CTE), and high thermal conductivity relative to austenitic stainless steel (Ref 2). During the welding of martensitic stainless steels, the fusion zone and that portion of the HAZ that is heated to a temperature higher than the austenitization temperature is transformed to a brittle, untempered martensite upon cooling to room temperature (Ref 4). Because the ductility, or fracture toughness, of the untempered martensite is inferior to that of the surrounding quenched-and-tempered microstructure, the material is unfit for service in the as-welded condition. Consequently, a postweld heat treatment (PWHT) must be employed to reduce the hardness and increase the toughness of the weld region. Recommended PWHT practices for martensitic stainless steels, including HT9, are described in the section "Specific Welding Recommendations" in this article. The as-welded properties of the weld fusion zone and its response to PWHT can often be controlled by the appropriate choice of weld filler metal (Ref 4). However, the properties of the HAZ can be controlled only by the judicious selection of the welding process and process parameters, and by the use of a PWHT. As a result, the utilization of martensitic

stainless steels depends, to a large extent, on the ability of the user to understand and control the properties of the weld HAZ. An extensive review of the microstructure and properties of the HAZ in martensitic HT9 weldments is provided by Lippold (Ref 2, 3, 4). In his studies, a variety of autogenous bead-on-plate gas-tungsten arc welds were evaluated using both light optical microscopy and microhardness techniques. Prior to welding, the base-metal plates were austenitized for 30 min at 1040 °C (1900 °F), followed by an air cool. Tempering was performed at 760 °C (1400 °F) for 1 h. The basemetal microstructure consisted of a mixture of tempered lath martensite and alloy carbides, as shown in Fig. 2.

FIG. 2 QUENCHED-AND-TEMPERED BASE METAL (HT9) MICROSTRUCTURE, 22 TO 25 HRC. SOURCE: REF 2

The welding parameters used in these studies are listed in Table 5. The welding current was varied in the range from 100 to 225 A in order to produce weld heat inputs ranging from 0.71 to 2.13 kJ/mm (18 to 54 kJ/in.). The plates were clamped in a copper fixture during welding in order to stimulate the structural restraint experienced by actual components during fabrication.

TABLE 5 PARAMETERS FOR GAS-TUNGSTEN ARC WELDING OF HT9 MARTENSITIC STAINLESS STEEL

VOLTAGE, V ARC CURRENT, A TRAVEL SPEED, MM/S (IN./MIN) ELECTRODE/WORK DISTANCE, MM (IN.) ELECTRODE DIAMETER, MM (IN.) ELECTRODE MATERIAL SHIELDING GAS (ARGON), MM3/S (FT3/H) PREHEAT

12-16 100-225 1.69 (4) 2.38 (0.094) 3.18 (0.125) W-2THO2 365 (30) NONE

Source: Ref 2, 3

As shown in Fig. 3, the weld area consisted of the base metal (described above and shown in Fig. 2), the HAZ, and the fusion zone. Prior to tempering, the fusion zone microstructure comprised a mixture of untempered martensite and ferrite, as shown in Fig. 4.

FIG. 3 FOUR HAZ REGIONS OBSERVED IN HT9 WELDMENTS. γ REPRESENTS AUSTENITE. SOURCE: REF 3

FIG. 4 AS-WELDED GAS-TUNGSTEN ARC FUSION ZONE MICROSTRUCTURE IN HT9 WELDMENTS. ARROWS INDICATE METASTABLE δFERRITE, WHICH CONSTITUTES APPROXIMATELY 2 TO 3 VOL% OF THE STRUCTURE. 1000×. SOURCE: REF 2

Four distinct regions were identified within the HAZ. These regions can be related to the phase fields on the ironchromium-carbon equilibrium diagram shown in Fig. 5. The microstructures associated with these four regions are (Ref 4): •

•

•

•

REGION 1, WHICH IS A TWO-PHASE REGION ADJACENT TO THE FUSION LINE THAT CONSISTS OF UNTEMPERED MARTENSITE AND FERRITE. THE FERRITE FORMS ALONG PRIOR AUSTENITE GRAIN BOUNDARIES AS THIS REGION IS HEATED INTO THE AUSTENITE-PLUS-FERRITE PHASE FIELD (DESIGNATED 1 IN FIG. 3 AND 5). REGION 2, WHICH CONSISTS ENTIRELY OF UNTEMPERED MARTENSITE. AS THIS PORTION OF THE HAZ IS HEATED INTO THE HIGH-TEMPERATURE REGIME OF THE AUSTENITE PHASE FIELD, ALLOY CARBIDES FROM THE ORIGINAL BASE-METAL MICROSTRUCTURE COMPLETELY DISSOLVE, RESULTING IN A CARBIDE-FREE UNTEMPERED MARTENSITE THAT EXHIBITS A LARGE PRIOR AUSTENITE GRAIN. REGION 3, WHICH CONTAINS UNTEMPERED MARTENSITE INTERSPERSED WITH UNDISSOLVED ALLOY CARBIDES. BECAUSE THIS REGION EXPERIENCES A TEMPERATURE EXCURSION INTO THE LOWER REGIME OF THE AUSTENITE PHASE FIELD, CARBIDE DISSOLUTION IS INCOMPLETE AND THE PRIOR AUSTENITE GRAIN SIZE IS SMALLER RELATIVE TO REGION 2. REGION 4, WHICH CONSISTS OF TEMPERED MARTENSITE AND OVERAGED CARBIDES. IT APPEARS TO BE VIRTUALLY IDENTICAL TO THE BASE-METAL MICROSTRUCTURE.

Representative microstructures of these four regions are shown in Fig. 6.

FIG. 5 IRON-CHROMIUM-CARBON PSEUDO-BINARY PHASE DIAGRAM FOR 12 WT% CR STEEL. CIRCLED NUMBERS REPRESENT THE FOUR HAZ REGIONS SHOWN IN FIG. 3 AND 6. SOURCE: ADAPTED FROM REF 5

FIG. 6 GAS-TUNGSTEN ARC WELDED HAZ MICROSTRUCTURES IN HT9 WELDMENTS. SOURCE: REF 2

The microstructures in the HAZ can also be differentiated by means of a microhardness traverse, as shown in Fig. 7. The drop in hardness in Region 1 results from the presence of the softer ferrite phase situated at prior austenite grain boundaries within the untempered martensite matrix. The highest hardness in the HAZ occurs in the region where carbide dissolution is nearly complete. Because the hardness of the martensite is almost entirely a function of the carbon content of the prior austentite, Region 2 exhibits the peak HAZ hardness. Within Region 3, the carbon content of the austenite is

less, because of the incomplete dissolution of the carbides at lower austenitization temperatures. Therefore, the resultant microstructure, which consists of untempered martensite and residual carbides, is softer than the adjacent untempered microstructure (Region 2). The hardness decrease in Region 4 results from overtempering the base-metal microstructure at a temperature just below Ac1 (the temperature at which austenite begins to form during heating), which is approximately 840 °C (1545 °F).

FIG. 7 MICROHARDNESS TRAVERSE ACROSS A GAS-TUNGSTEN ARC WELD IN HT9 STAINLESS STEEL (NO PREHEAT OR POSTHEAT TREATMENT). SOURCE: REF 3

Figures 8 and 9 demonstrate the effect of PWHT on the hardness and toughness, respectively, of the four HAZ regions described above. As shown in these figures, a PWHT of 760 °C (1400 °F) for 1 h produced superior results.

FIG. 8 HARDNESS AS A FUNCTION OF PWHT TEMPERATURE IN THE FOUR HAZ REGIONS OF HT9 WELDMENTS. SOURCE: REF 4

FIG. 9 EFFECT OF ORIENTATION AND PWHT ON THE TOUGHNESS OF HT9 WELDMENTS. (A) REGION 1. (B) REGION 2. (C) REGION 3. (D) REGION 4. SEE ALSO FIG. 3 AND 6. SOURCE: REF 4

Similar studies on HT9 have been carried out by Fenn and Jordan, with particular attention being given to δ ferrite formation (Ref 6). Maintaining a δ ferrite content that is as low as possible (< 10 vol%) is important, because a homogeneous martensitic structure that is free of δ ferrite islands results in a tougher weld. The investigations of these workers covered a variety of welding processes and a wide range of welding heat inputs, and involved real and simulated HAZ material.


2. J.C. LIPPOLD, TRANSFORMATION AND TEMPERING BEHAVIOR OF 12CR-1MO-0.3V MARTENSITIC STAINLESS STEEL WELDMENTS, J. NUCL. MATER., VOL 103 AND 104, 1981, P 1127-1132 3. J.C. LIPPOLD, "MICROSTRUCTURAL CHARACTERIZATION OF AUTOGENOUS GTA WELDS IN A 12CR-1MO-0.3V STEEL," DOCUMENT NO. SAND80-8236, SANDIA NATIONAL LABORATORIES, OCT 1980 4. J.C. LIPPOLD, THE EFFECT OF POSTWELD HEAT TREATMENT ON THE MICROSTRUCTURE AND PROPERTIES OF THE HEAT-AFFECTED ZONE IN 12CR-1MO-0.3V (HT9) WELDMENTS, PROC. TOPICAL CONFERENCE ON FERRITIC ALLOYS FOR USE IN NUCLEAR ENERGY TECHNOLOGIES, TMS, 1984, P 497-506 5. J.Z. BRIGGS AND T.D. PARKER, THE SUPER 12% CR STEELS, CLIMAX MOLYBDENUM

COMPANY, 1965 6. R. FENN AND M.F. JORDAN, MICROSTRUCTURE OF WELD HEAT-AFFECTED ZONE IN 12CR1MO STEEL, MET. TECHNOL., VOL 9 (NO. 8), 1982, P 327-337 Selection of Wrought Martensitic Stainless Steels J.R. Davis, Davis & Associates

Weldability Martensitic stainless steels can be welded in the annealed, hardened, and hardened-and-tempered conditions. Regardless of the prior condition of the steel, welding produces a hardened martensitic zone adjacent to the weld. In other words, the high-temperature HAZ will be in the "as-quenched" condition after welding, regardless of the prior condition of the material. In addition, the HAZ hardness is very much independent of the cooling rate over the temperature range experienced in common arc welding practices. This is evident from the isothermal transformation diagram shown in Fig. 10. Because such high hardness values render the material prone to cracking during fabrication, the selection of appropriate preheating levels and welding procedures is critical to the success of the welding process.

FIG. 10 ISOTHERMAL TRANSFORMATION DIAGRAM FOR TYPE 410 STAINLESS STEEL

Hydrogen-Induced Cold Cracking. Weld-area cracking in martensitic stainless steels is primarily due to the presence of hydrogen in the hardened structure (Ref 7). Hydrogen-induced cold cracks form in welds when the weldment is at or

near room temperature (typically, less than 150 °C, or 300 °F). Cracking may occur almost immediately or hours after cooling. In steel, cold cracks depend on the presence of a tensile stress, a susceptible microstructure, and sufficient hydrogen in the weld. Elimination of one or more of these factors greatly reduces crack susceptibility. The stress may arise from restraint by other components of a weldment or from the simple thermal stresses created by welding in a butt, groove, or "T" joint. The susceptibility of the microstructure to cold cracking relates to the solubility of hydrogen and the possibility of supersaturation. Austenite, in which hydrogen is highly soluble, is least susceptible to cold cracking, and martensite, in which the solubility is lower, is most susceptible, because the rapid cooling necessary for the austenite-to-martensite transformation traps the hydrogen in a state of supersaturation in the martensite. Because the solubility of hydrogen in bcc martensite is low and diffusivity is high, austenitic filler metals are often selected to weld martensitic stainless steels. The discussion on filler metal selection later in this article provides more details. The presence of hydrogen in a weld is generally due to moisture that is introduced in the shielding gas (or the electrode coating or flux), dissociated by the arc to form elemental hydrogen, and dissolved by the molten weld pool and by the adjacent region in the HAZ. In the supersaturated state, the hydrogen diffuses to regions of high stress, where it can initiate a crack. Continued diffusion of the hydrogen to the region of stress concentration at the crack tip extends the crack. This behavior means that hydrogen-induced cold cracking is time dependent--that is, time is needed for hydrogen diffusion--and the appearance of detectable cracks can be delayed until long after the weld has passed inspection. Additional information on the mechanisms of hydrogen cracking of weldments can be found in Ref 8 and 9. For most commercial steels, the avoidance of hydrogen-induced cold cracking is based on the control of hydrogen (Ref 7). The sources of hydrogen (water, oils, greases, waxes, and rust that contains hydrogen or hydrates) should be eliminated. The hydrogen potential also can be minimized by using low-hydrogen, inert-gas welding processes, such as gas-tungsten arc welding (GTAW) or gas-metal arc welding (GMAW), or by paying stringent attention to consumable drying and baking for flux-shielded welding processes, such as shielded metal arc welding (SMAW) or flux-cored arc welding (FCAW). In addition, preheating must be applied to slow the rate of cooling. This allows more time for hydrogen to diffuse away from the weld area during cooling within the austenite range and, especially, following transformation to martensite. Preheating is generally carried out in the temperature range from 200 to 300 °C (400 to 600 °F). In multipass welds, interpass temperatures must be maintained at the same level, and it is frequently beneficial to hold this temperature for some time after arc extinction to permit further hydrogen diffusion out of the joint. Postweld tempering is carried out to improve weld-area toughness. For higher-carbon steels (>0.2 wt% C), this process should be applied as soon as possible after welding to avoid the possibility of hydrogen-induced cracking from atmospheric corrosion. The weld must be cooled to a sufficiently low temperature to induce the austenite-to-martensite transformation (in principle, below the martensite completion point, Mf, as described below) prior to PWHT. Material Composition and Selection of Preheat Temperature (Ref 10). The temperature at which austenite

transforms to martensite upon cooling in martensitic stainless steels depends on material composition, particularly the carbon, manganese, nickel, chromium, and molybdenum contents. The martensitic start temperature, Ms, for such steels can be calculated using the following equation developed by The Welding Institute (Ref 11):

MS = 540 - (497 × %C + 6.3 × %MN + 36.3 × %NI + 10.8 × %CR + 46.6 × %MO) °C

(EQ 1)

It should be noted that all of the elements listed above lower the Ms temperature. Using Eq 1 for a type 410 stainless steel of the composition listed in Table 1, the Ms temperature is approximately 350 °C (660 °F). The temperature at which the transformation to martensite is complete, denoted by Mf, is approximately 100 °C (180 °F) below the Ms, that is, 256 °C (493 °F) for type 410. The preheat level should be below the Mf wherever possible, for example, 250 °C (480 °F) for type 410. However, if the Mf is less than about 150 °C (300 °F), as is the case with highcarbon (>0.2 wt% C) steels, it is not advisable to select a lower preheat level, because the embrittling effect of hydrogen can be significant at this temperature. In most cases, the reason for choosing a preheat temperature below the Mf is to ensure that maximum transformation to martensite occurs. If this has not happened, then subsequent PWHT immediately after welding may be ineffective for two

reasons. First, the partially austenitic structure will retain hydrogen, which can lead to hydrogen cracking problems upon eventual cooling to room temperature, where untempered martensite with a high hydrogen concentration could form. Second, an isothermal transformation of the weld can take place during heat treatment, and the martensite that forms upon cooling to room temperature will be untempered and hard. The major compositional factor influencing hydrogen-induced cold cracking is the material carbon content, because it determines the hardness of the transformed structure and its sensitivity to hydrogen embrittlement (Ref 7). Below levels of approximately 0.1 wt% C, and at least in thin sections (41, DUPLEX STAINLESS STEELS CONF. PROC., VOL 2, LES EDITIONS DE PHYSIQUE, LES ULIS CEDEX, FRANCE, OCT 1991, P 1341-1351 12. H. ERIKSSON AND S. BERNHARDSSON, THE APPLICABILITY OF DUPLEX STAINLESS STEELS IN SOUR ENVIRONMENTS, CORROSION, VOL 47 (NO. 8), SEPT 1991, P 719-727 13. A. MLYASAKA AND H. OGAWA, CRITICAL STRESS FOR SCC OF CORROSION RESISTANT ALLOY IN SOUR ENVIRONMENTS, PAPER 5, CORROSION '89, NATIONAL ASSOCIATION OF CORROSION ENGINEERS, 1989 14. S. BERNHARDSSON, THE CORROSION RESISTANCE OF DUPLEX STAINLESS STEELS, DUPLEX STAINLESS STEELS CONF. PROC., AMERICAN SOCIETY FOR METALS, 1983, P 185-210 15. S. AZUMA, H. TSUGE, T. KUDO, AND T. MOROISHI, CREVICE CORROSION OF DUPLEX STAINLESS STEELS IN SIMULATED SOUR GAS ENVIRONMENTS, PAPER 306, CORROSION '87, NATIONAL ASSOCIATION OF CORROSION ENGINEERS, 1987 16. "SULFIDE STRESS CRACKING RESISTANT METALLIC MATERIALS FOR OILFIELD EQUIPMENT," MR-01-75-93, NATIONAL ASSOCIATION OF CORROSION ENGINEERS 17. L. SMITH, CORROSION RESISTANT ALLOYS OFFER PIPELINE ALTERNATIVES, STAINL. STEEL EUR., JAN/FEB 1993, P 32-42 18. M.J. SCHOFIELD, AND R.D. KANE, DEFINING SAFE USE LIMITS FOR DUPLEX STAINLESS STEELS, DUPLEX STAINLESS STEEL CONF. PROC., LES EDITIONS DE PHYSIQUE, LES ULIS CEDEX, FRANCE, OCT 1991, P 241-255

19. T. OGAWA AND T. KOSEKI, EFFECT OF COMPOSITION PROFILES ON METALLURGY AND CORROSION BEHAVIOUR OF DUPLEX STAINLESS STEEL WELD METALS, WELD. J., VOL 88 (NO. 5), MAY 1989, P 181S-191S Selection of Wrought Duplex Stainless Steels David N. Noble, ARCO Alaska, Inc.

Applications Duplex stainless steels have found widespread use in a range of industries, particularly the oil and gas, petrochemical, pulp and paper, and pollution control industries. They are commonly used in aqueous, chloride-containing environments and as replacements for austenitic stainless steels that have suffered from either Cl SCC or pitting during service (Ref 20). The more highly alloyed grades (that is, the "super" DSS) are resistant to oxygenated or chlorinated sea water (Ref 21). The leaner alloy materials are not resistant to live sea water and will suffer from pitting. Several grades have been used for transportation of wet, CO2-containing gas where carbon would have suffered general corrosion (Ref 22). Table 6 gives a list of DSS grades and typical applications.

TABLE 6 TYPICAL APPLICATIONS OF DUPLEX STAINLESS STEELS

GRADE(S) APPLICATION S31803, S32760, TUBING FOR HEAT EXCHANGERS IN REFINERIES, CHEMICAL INDUSTRIES, S32750, S32550 PROCESS INDUSTRIES, AND OTHER INDUSTRIES USING WATER AS A COOLANT; KRAFT PAPER DIGESTERS ALL HEAT EXCHANGERS, CHEMICAL TANKERS, CHEMICAL REACTOR VESSELS, FLUE GAS FILTERS, ACETIC AND PHOSPHORIC ACID HANDLING SYSTEMS, OIL AND GAS INDUSTRY EQUIPMENT (MULTIPHASE FLOW LINES, DOWNHOLE PRODUCTION TUBULARS, COMMONLY COLD WORKED) S32304 DOMESTIC WATER HEATERS AND WHERE PITTING RESISTANCE IS NOT OF OVERRIDING IMPORTANCE S32750, S32760, PIPE FOR SEAWATER HANDLING AND FIREFIGHTING SYSTEMS, OIL AND S32550 GAS SEPARATORS, SALT EVAPORATION EQUIPMENT, DESALINATION PLANTS, GEOTHERMAL WELL HEAT EXCHANGERS, HUMAN BODY IMPLANTS (S32550 MAY SUFFER SLIGHT PITTING AND CREVICE CORROSION IN SEAWATER SERVICE) References cited in this section

20. D. FRUYTIER, INDUSTRIAL EXPERIENCES WITH DUPLEX STAINLESS STEEL, STAINL. STEEL EUR., VOL 3 (NO. 13), DEC 1991 21. B. WALLEN AND S. HENRIKSON, EFFECT OF CHLORINATION ON STAINLESS STEELS IN SEAWATER, PAPER 403, CORROSION '86, NATIONAL ASSOCIATION OF CORROSION ENGINEERS, 1986 22. J. HARSTON, E. HUTCHINS, AND S. SWEENEY, THE DEVELOPMENT AND CONSTRUCTION OF DUPLEX STAINLESS STEEL PIPELINES FOR USE OFFSHORE IN THE SOUTHERN NORTH SEA, PROC. 3RD LNT. CONF. WELDING AND PERFORMANCE OF PIPELINES, THE WELDING INSTITUTE, 1986

Selection of Wrought Duplex Stainless Steels

David N. Noble, ARCO Alaska, Inc.

Microstructural Development Factors Controlling Microstructure. By assuming a duplex stainless steel contains 68% Fe and 32% alloying elements (comprising nickel and chromium, but representing austenite and ferrite formers, respectively), the 68% Fe pseudo-binary phase diagram can be used to explain the microstructural transformations that occur during heating and cooling (Fig. 9).

FIG. 9 PSEUDO-BINARY DIAGRAM AT CONSTANT IRON SECTION (68 WT% FE) OF THE FE/CR/NI TERNARY SYSTEM. SEE TEXT FOR DEFINITION OF LINES A, B, AND C AND POINTS D AND E. SOURCE: REF 9

Composition. With approximately 30% Cr and 2% Ni (composition indicated by line A), the steel will solidify

completely to ferrite and remain ferritic down to room temperature. As the chromium content decreases and the nickel content increases, an alloy containing 28% Cr and 4% Ni, indicated by line B, will solidify as ferrite, remain fully ferritic until about 1200 to 1300 °C (2190 to 2370 °F), and then enter the dual-phase field (denoted by α + γ). A two-phase microstructure will be developed and maintained down to room temperature. Because the composition of alloy B lies to the far left of the dual-phase field, a microstructure rich in ferrite would be anticipated. Current commercial DSS with 25% Cr are more likely to contain about 7% Ni, indicated by line C, which upon cooling enters the dual-phase field at a much higher temperature, thereby allowing more time for the diffusion-driven solid-state transformation from ferrite to austenite. An alloy with composition C will solidify as ferrite and be fully ferritic for a narrow temperature range prior to entering the dual-phase field. It should be noted that alloys A, B, and C will solidify primarily as ferrite and will have ferritic matrices. Heat Treatment. The sequence of phase transformation in DSS upon cooling is (Fig. 9):

LIQUID → LIQUID + FERRITE → FERRITE → FERRITE + AUSTENITE

(EQ 2)

The solid-state transformation from ferrite to austenite is considered to be diffusional (Ref 23) and therefore affected by time and temperature. Very rapid cooling from a temperature at which the steel is fully ferritic will result in suppression

of the ferrite-to-austenite transformation and retention of a predominantly ferritic structure. The temperature range over which the transformation from ferrite to austenite occurs is highly dependent upon composition and grade (Fig. 9), but is typically from 1200 to 800 °C (2190 to 1470 °F). For this reason, a ∆T12-8 cooling rate term has been used when characterizing a DSS thermal cycle (Ref 24, 25). A longer ∆T12-8 will result in more transformation from ferrite to austenite. In a weld of given thickness, longer ∆T12-8 values can be achieved by a higher heat input weld. The ∆T12-8 term has been employed in an effort to identify critical cooling rates necessary to develop acceptable austenite content in weld heat-affected zones to achieve the desired mechanical and corrosion resistance properties. The critical ∆T12-8 for phase transformation is independent of material thickness, unlike heat input, but is highly dependent on alloy composition. Low-temperature thermal cycles will not affect the austenite/ferrite balance, but they will control the extent of intermetallic phase precipitation. These steels should undergo rapid cooling (that is, water quenching) through the precipitation temperature range to develop optimum properties. For welds, water quenching is usually not necessary or practical, but some means of forced cooling may be beneficial for productivity and quality reasons, particularly for multipass welds in the high-alloy grades (Ref 26, 27). To produce the necessary base material microstructure, an isothermal heat treatment is carried out at a temperature at which the steel is dual phase, approximately 1050 to 1100 °C (1920 to 2010 °F), or point D in Fig. 9 (Ref 28). To avoid the formation of intermetallic phases during cooling, water quenching from the soaking temperature is necessary. Morphology. Wrought DSS base material microstructures have a pronounced orientation of austenite islands in the

ferritic matrix, parallel and transverse to the rolling direction, as a result of the hot and cold working following casting (Fig. 1). The cast microstructure (Fig. 2) is very different, and more similar to that observed in weld metal and the HAZ. During the weld thermal cycle, the HAZ adjacent to the fusion line will become fully ferritic at temperatures above the austenite-to-ferrite transition temperature (point E in Fig. 9). Grain growth will occur and, regardless of the extent of austenite transformation that occurs during subsequent cooling, the coarse prior ferrite grain boundaries will still be apparent (Fig. 10). Very rapid cooling from the fully ferritic phase field will retard austenite reformation in the HAZ, and result in a predominantly ferritic structure, also shown in Fig. 10. When cooled more slowly, the austenite will form initially on the ferrite grain boundaries and then within the ferrite grains along preferred crystallographic planes in a manner similar to the cast material.

FIG. 10 PHOTOMICROGRAPH OF A FE-22CR-5.5NI-3MO-0.15N DSS WELD AREA SHOWING ELONGATED PARENT PLATE MORPHOLOGY, A COARSE GRAIN HEAT-AFFECTED ZONE HIGH IN FERRITE, AND COLUMNAR WELD METAL GRAINS. SPECIMEN ELECTROLYTICALLY ETCHED IN 40% NAOH. 30×. COURTESY OF THE WELDING INSTITUTE

The HAZ in Fig. 10 is relatively wide and has a high ferrite content. Such a structure may lead to poor low-temperature notch toughness and reduced corrosion resistance. Nitrogen levels in the base steel are being steadily increased to improve weldability and to avoid the formation of structures similar to that shown in Fig. 10; some DSS contain as much as 0.35 wt% N (Ref 4). A higher nitrogen level will increase the temperature at which ferrite begins to transform to austenite on cooling (Ref 6). This reduces the width of the transformed HAZ by limiting to a narrow temperature range the extent of the fully ferritic phase field. It also encourages more austenite reformation by starting the transformation at a much higher temperature. The recent grades of DSS, which typically have minimum 0.20 wt% N, develop a very narrow HAZ in which the transformed region is often difficult to distinguish. The sequence of transformation reactions in the weld metal is the same as for the HAZ, but due to the solidification pattern, the macroappearance is different and columnar grains are more typical (this can be seen on the top left of Fig.

10). Duplex weld metal grain growth is epitaxial, with the large prior ferrite grains in the HAZ dictating weld metal grain size and orientation. Austenite again forms preferentially on the grain boundaries, and intragranular austenite forms along preferred crystallographic planes as the weld metal continues to cool (Fig. 11).

FIG. 11 PHOTOMICROGRAPH OF DSS WELD METAL WITH AUSTENITE LATHS (WHITE) AT PRIOR FERRITE GRAIN BOUNDARIES AND ALONG PREFERRED CRYSTALLOGRAPHIC PLANES IN A FERRITIC MATRIX. SAMPLE ELECTROLYTICALLY ETCHED IN 40% NAOH. 100×. COURTESY OF THE WELDING INSTITUTE

Precipitates are often difficult to observe and require precise etching techniques and experience with the type of microstructure. Due to its highly detrimental effect on corrosion resistance and toughness and relative ease of formation, the most commonly investigated precipitate in DSS is σ phase, an Fe(Cr,Mo) intermetallic. Sigma will initially precipitate at phase boundaries in the ferrite phase and will then grow into the ferrite phase, which is richer in chromium and molybdenum. Only a small amount of σ precipitation is required to reduce mechanical and corrosion properties.


4. L. VAN NASSAU, H. MEELKER, AND J. HILKES, WELDING DUPLEX AND SUPER-DUPLEX STAINLESS STEELS--A GUIDE FOR INDUSTRY, DOCUMENT 01463, LINCOLN NORWELD, JULY 1992 6. J. CHARLES, SUPER DUPLEX STAINLESS STEEL: STRUCTURE AND PROPERTIES, DUPLEX STAINLESS STEELS CONF. PROC., VOL 1, LES EDITIONS DE PHYSIQUE, LES ULIS CEDEX, FRANCE, OCT 1991, P 3-48 9. R.J. BRIGHAM AND E.W. TOZER, EFFECT OF ALLOYING ADDITIONS ON THE PITTING RESISTANCE OF 18% CR AUSTENITIC STAINLESS STEEL, CORROSION, VOL 30 (NO. 5), MAY 1974, P 161-166 23. J.E. LIPPOLD AND W.A. BAESLACK, MET. CONSTR., VOL 20 (NO. L), 1988, P 21 24. H. HOFFMEISTER AND R. MUNDT, ARCH. EISENHUTT., VOL 52 (NO. 4), P 159-164 25. D.N. NOBLE AND T.G. GOOCH, THE EFFECTS OF HEAT INPUT AND CHEMICAL COMPOSITION ON MICROSTRUCTURE IN ARC WELDED UNS S31803 TYPE DUPLEX STAINLESS STEELS, ASM MATERIALS WEEK '87, ASM INTERNATIONAL, OCT 1987 26. B. JOSEFSSON, J.-O. NILSSON, AND A. WILSON, PHASE TRANSFORMATION IN DUPLEX STEELS AND THE RELATION BETWEEN CONTINUOUS COOLING AND ISOTHERMAL HEAT TREATMENTS, DUPLEX STAINLESS STEELS CONF. PROC., AMERICAN SOCIETY FOR METALS, 1983 27. L. ITRURGOYEN, J. ALCALA, AND M. ANGLADA, THE INFLUENCE OF AGEING AT 475 C ON THE FRACTURE RESISTANCE OF A DUPLEX STAINLESS STEEL, CORROSION, VOL 47 (NO. 8), SEPT 1991, P 757-764 28. "SANDVIK SAF 2507--A HIGH PERFORMANCE DUPLEX STAINLESS STEEL," SANDVIK STEEL TRADE LITERATURE, PAMPHLET S-1875-ENG, SEPT 1987

Selection of Wrought Duplex Stainless Steels David N. Noble, ARCO Alaska, Inc.

Welding and Weldability Duplex stainless steel weldability is generally good, although it are not as forgiving as austenitic stainless steels or as prone to degradation of properties as fully ferritic stainless steels. The current commercial grades are low in carbon (less than 0.03 wt%), thereby essentially eliminating the risk of sensitization and intergranular corrosion from carbide precipitation. The base material and filler metals also have low sulfur and phosphorus levels (less than 0.03 wt%), which in combination with the ferritic solidification reduce the likelihood of solidification cracking (hot cracking). Hydrogen cracking (cold cracking) resistance is also good due to the high hydrogen solubility in the austenite (Ref 29) and the high percentage of austenite in the matrix. Nevertheless, solidification cracking and hydrogen cracking can occur in duplex alloys, and some precautions are necessary. Preheat is generally not recommended for DSS, but may sometimes be specified in low-nitrogen grades, because thick sections and low heat input welding processes may, in combination, develop highly ferritic heat-affected zones (Ref 30). For the more highly alloyed DSS, a preheat can be highly detrimental and reduce corrosion resistance and mechanical properties. Postweld heat treatment (PWHT) is not commonly used except in autogenous welds or welds with a filler metal composition that exactly matches the base steel. Although not always necessary, particularly if a nickel-enriched filler metal is used, it is common to postweld heat treat DSS welded pipe after longitudinal seam welding. The PWHT will largely be for the purpose of restoring the correct phase balance and redissolving unwanted precipitates. Postweld heat treatment temperatures of about 1050 to 1100 °C (1920 to 2010 °F) are used, depending on grade, followed by the same heat treatment applied to the base material during solution annealing--usually water quenching (Ref 31). The heat treatments commonly used for structural steels (for example, 550 to 600 °C, or 1020 to 1110 °F) (Ref 32) are totally inappropriate for duplex alloys and should never be considered. Interpass Temperature Control. The need for interpass temperature control depends on alloy grade and composition.

The more highly alloyed grades should have tighter interpass control. For the Fe-22Cr-5.5Ni-3Mo-0.15N steels, interpass temperatures of 150 to 200 °C (300 to 390 °F) are typical, whereas for the Fe-25Cr-7Ni-3.5Mo-0.25N-(Cu-W) steels, temperatures as low as 70 °C (160 °F) have been recommended (Ref 33). Welding practices employed for austenitic stainless steels (for example, cleanliness, use of a backing gas, avoidance of

contamination with carbon steel, and so on) should also be adopted for DSS. The backing gas most commonly used is pure argon. However, the use of a backing gas without nitrogen can lead to nitrogen loss from the molten weld nugget (Ref 34) with probable loss in corrosion performance and toughness. Because of this, argon/nitrogen mixed gases for shielding and backing have been explored (Ref 35), and generally the corrosion resistance of welds produced with these gases has been superior to that of welds made with nitrogen-free gases. Pure nitrogen may also be used as a backing gas (Ref 2), but a concern exists that too much nitrogen may be picked up, thereby significantly affecting the phase balance locally by promoting the formation of austenite. Open root gaps are commonly used in manual DSS welding to limit the extent of dilution from the base material, because extensive base metal dilution will negate the effect of the nickel over-alloying in the filler metal. To control the problems of too little austenite and intermetallic phase formation, associated with too-low and too-high heat inputs, respectively, DSS welding procedures typically specify a "window" of allowable heat inputs. For the highly alloyed and heat-sensitive super DSS, stringer beads are commonly specified, particularly in the hot pass (that is, weld pass immediately following the root), and in some cases in split hot passes (that is, two hot passes side by side). The technique of using a low heat input hot pass has led to the use of the term "cold" pass (Ref 33). Welding procedure qualification generally follows the mechanical property test requirements of a major code, for

example, ASME Section IX (Ref 36) or API 1104 (Ref 37). There are also several tests (specifically phase balance assessment, microstructural examination, and a pitting corrosion test) used for DSS weldment evaluation that are not typically employed for other stainless steels.

Phase-Balance Assessment. The phase balance test measures the amount of austenite and ferrite present in the

microstructure and is similar to tests used to measure ferrite in austenitic weld metals. However, the testing requirements are usually more exhaustive for duplex alloys and will typically call for assessment of the ferrite content in the HAZ and weld metal. and possibly at several through-thickness locations (Ref 38). The amount of ferrite in a duplex structure can be specified and measured in two ways. Percentage ferrite is measured by manual point counting or image analysis. These methods measure the area fraction of ferrite and austenite exposed on the surface of a sample under examination. The values are reported as a volume fraction, thereby linking the surface area measurement on one plane to the overall volume fraction in the sample. Percentage ferrite requirements are typically 35 to 65% and are applied to weld metal and the HAZ. The alternative method uses the ferromagnetic properties of the ferrite to measure its proportion in the matrix. The Ferrite Number (FN) system used for austenitic stainless steel weld metals (where the ferrite content is measured by attraction of a spring-loaded magnet) was not originally developed for high ferrite levels, such as those in duplex alloys, and was only calibrated to about 28 FN (Ref 39). Furthermore, the variations in iron content between the various duplex grades affected the magnetic attraction for a given phase balance (Ref 40). Because no reliable magnetic attraction system had been fully developed and proven at the time duplex alloys were seeing widespread industrial application, most specifications required ferrite percentage measurements. However, changes to the FN measuring system have now been made so that all ferrite levels in duplex alloys can be recorded, and the system has been shown to be more reliable. To achieve the extended range, counterweights were added to the spring-loaded magnet (Ref 40). Specification requirements for weld metals are typically 50 to 90 FN. A rule-of-thumb conversion from FN to percentage is:

0.70(FN) = % FERRITE

(EQ 3)

Measurement of ferrite content in DSS, however, is seldom currently performed with a magnetic measuring technique and is more commonly referred to in terms of percentage. There are also some limitations to the application of the FN measuring system in certain regions of the weld. The measurement tool applies well to large areas of finely dispersed ferrite, such as in the weld metal, but it cannot accurately measure small areas that may have locally high ferrite levels (for example, the weld HAZ). It also may provide inaccurate readings in base material measurements due to the effect of magnet deflection toward the relatively coarse base material ferritic matrix. Microstructural Examination. Optical microstructural examination is included to determine whether extensive phase

precipitation has occurred (for example, σ-phase, carbides, nitrides, and so on), and specifications frequently call for the absence of any deleterious phases. However, the interpretation of DSS microstructures is so difficult that one cannot make a definitive assessment, especially without experience in studying similar structures. Detecting small proportions of precipitates by optical means is also relatively difficult. Pitting Corrosion Test. The pitting corrosion test most commonly used for qualifying DSS is a modification of the

ASTM G 48 Practice A ferric chloride test (Ref 41). Other tests are available (for example, potentiostatic critical pitting temperature tests in 3 or 15% NaCl, and others involving the use of CO2), but the ASTM G 48 test has been shown to detect material with substandard properties and is preferred to most other common test techniques (Ref 42). The test solution is not representative of most service conditions in that it is highly acidic, oxidizing, and has a high chloride concentration, but the test is relatively easy to perform and can distinguish between good and bad welds. As used on DSS, the test involves the immersion of a welded coupon into a 10% FeCl3 · 6H2O solution for a duration of between 24 and 72 h at a temperature from 15 to 70 °C (60 to 160 °F), depending on material grade and specification requirements. Test assessment can be by direct weight loss, visual observation of pitting, weight loss per unit of exposed area, or corrosion rate. The procedure as laid out by ASTM was not designed for DSS weldment testing, and although it has been adapted in a number of ways, its interpretation still presents difficulties (Ref 43). Filler Metal Requirements. For most DSS grades there are two types of filler metals: • •

MATCHING COMPOSITION FILLER METALS FILLER METALS THAT ARE SLIGHTLY OVERALLOYED, PRINCIPALLY WITH RESPECT TO NICKEL

The matching filler metal is used where a PWHT is performed, whereas welds made with the filler metal enriched with nickel are used in the as-welded condition. The weld metal microstructure from a composition exactly matching that of the parent steel will contain a high ferrite content. The increase in nickel is made to improve the as-welded phase balance and increase austenite content. The ferrite content of a weld made with a nickel-enriched consumable would decrease significantly if it underwent a PWHT. It may suffer from slightly reduced weld metal strength and could also be more susceptible to σ phase formation during heat treatment (Ref 44). The nickel level in the enriched weld metal will be approximately 2.5 to 3.5% greater than in the base material (for example, for the Fe-22Cr-5.5Ni-3Mo-0.15N DSS base material containing 5.5% Ni, the filler metal will comprise 8.0 to 9.0% Ni, depending on consumable manufacturer and form). Nominal filler metal compositions for the Fe-22Cr-5.5Ni-3Mo-0.15N grade of DSS, proprietary names, and the American Welding Society classifications, where appropriate, are reported in Table 7.

TABLE 7 TYPICAL COMPOSITION OF SELECTED COMMERCIAL NICKEL-OVERALLOYED WELDING CONSUMABLES AND AWS SPECIFICATIONS FOR FE-22CR-5.5NI-3MO-0.15N DUPLEX STAINLESS STEEL

MANUFACTURER

TRADE NAME COVERED ELECTRODES METRODE SUPERMET 2205 SANDVIK 22.9.3.L AVESTA 2205-PW LINCOLN NORWELD AROSTA 4462 THYSSEN THERMANIT 22/09 AWS A5.4 E2209 BARE WIRE SANDVIK 22.8.3.L METRODE ER329N THYSSEN 22/09/SG LINCOLN NORWELD LNT/LNM 4462 AWS A5.9(B) ER 2209 (A) (B)

COMPOSITION, % C Cr Ni

Mo

N

0.02 0.03 0.025 0.025 15

F

F

AUSTENITE + VERMICULAR δ FERRITE OR ACICULAR STRUCTURE WIDMANSTÄTTEN AUSTENITE IN FERRITIC MATRIX WIDMANSTÄTTEN AUSTENITE IN FERRITIC MATRIX WIDMANSTÄTTEN AUSTENITE IN FERRITIC MATRIX

(A) CREQ = %CR + 1.4(%MO) + 1.5(%SI) + 2(%NB) + 2.5(%AL) + 3(%TI); NIEQ = %NI + 0.3(%MN) + 35(%C) + 20(%N) + %CU. VALUES ARE FOR TYPICAL COMPOSITIONS. (B) A, AUSTENITIC; F, DELTA FERRITIC

FIG. 2 MICROSTRUCTURES OF AUTOGENOUS GTA WELD METAL IN THE AS-WELDED CONDITION. (A) 17-4PH, δ -FERRITE IN A MATRIX OF UNTEMPERED MARTENSITE. ETCHANT: FRY'S REAGENT. 440×. (B) PH13-8 MO, UNTEMPERED MARTENSITE. ETCHANT: VILELLA'S REAGENT. 440×. (C) 17-7PH, AUSTENITE PLUS ABOUT 20% δ -FERRITE. ETCHANT: ELECTROLYTIC NITRIC ACID-ACETIC ACID FOLLOWED BY ELECTROLYTIC OXALIC ACID. 440×. (D) A-286, FULLY AUSTENITIC STRUCTURE. ETCHANT: 90 ML HYDROCHLORIC ACID, 5 ML NITRIC ACID, 5 ML SULFURIC ACID, 1 ML HYDROFLUORIC ACID. 440×

In the as-welded condition, the heat-affected zones of welds in martensitic PH steels are distinguishable from the base metal by grain growth and an increased amount of ferrite in the high-temperature region close to the fusion line. Postweld solution treatments reduce the residual δ-ferrite content and refine the prior-austenite grain size of both the fusion and heat-affected zones (HAZ). Aging, likewise, simultaneously tempers the martensite and precipitates the agehardening particles in the fusion zone, HAZ, and base metal. However, these particles are too small to observe by optical microscopy. Aging may also reform up to 10% austenite at temperatures of 540 to 620 °C (1000 to 1150 °F), the amount increasing with aging temperature. 17-4PH Stainless Steel Weldability. 17-4PH stainless steel is typical of the martensitic PH stainless steels. Its weldability is superior to that of

the standard austenitic stainless steels because, due to its low carbon content, it does not require preheat to prevent cracking in thicknesses up to 100 mm (4.0 in.) (Ref 23 , 24, 25). The welding conditions used for arc welding 17-4PH stainless steel are essentially the same as those used for joining the standard austenitic stainless steels. However, neither welds nor base metal have the high ductility of the austenitic stainless steels, so care must be taken to avoid the presence of stress raisers (Ref 23 24, 25). One common form of stress raiser is the built-in notch at the root of a partial penetration weld. If design requirements dictate partial-penetration welds, the initiation of cracks at the root of the weld can be avoided by making the root pass with a ductile, low-strength filler metal such as type 308L, and then completing the balance of the weld with a matching high-strength heat-treatable filler metal. Arc Welding. Gas-tungsten arc, gas-metal arc, and shielded metal arc welding processes have all been used extensively

for joining 17-4PH stainless steel (Ref 23, 24, 25). Typical mechanical properties of arc welds in 17-4PH are presented in Table 6. The choice of welding process has little effect on the strength of welds after postweld heat treatment (PWHT),

but does affect the weld metal cleanliness and resultant ductility and toughness (Ref 23). Gas-tungsten arc welds have the highest and shielded metal arc welds the lowest quality. However, shielded metal arc welds are generally satisfactory for all but the most critical applications.

TABLE 6 MECHANICAL PROPERTIES OF 17-4PH ARC WELDS

THICKNESS POSTWELD HEAT (B)

TENSILE STRENGTH

ELONGATION, %

ksi

0.2% YIELD STRENGTH mpa ksi

MM

mpa

in 13 mm (0.5 in.)

in 25 mm (1 in.)

in 50 mm (2 in.)

AGED 480 °C (900 °F), 1 H AGED 540 °C (1000 °F), 4 H AGED 595 °C (1100 °F), 4 H AGED 480 °C (900 °F), 1 H AGED 540 °C (1000 °F), 4 H AGED 595 °C (1100 °F). 4 H ST + AGED 480 °C (900 °F), 1 H ST + AGED 510 °C (950 °F), 1 H ST + AGED 550 °C (1025 °F), 1 H

1340 1235 1140 1395 1240 1125 1370 1345 1215

194 179 165 202 180 163 199 195 176

1295 1170 1105 1290 1160 1060 1260 1255 1145

188 170 160 187 168 154 183 182 166

7 ... 14 16 17 ... ... ... ...

... ... ... ... ... ... 17.0 18.0 17.0

ST + AGED 480 °C (900 °F), 1 H ST + AGED 510 °C (950 °F), 1 H ST + AGED 550 °C (1025 °F), 1 H ST + AGED 480 °C (900 °F), 1 H

1400 1275 1195 1305

203 185 173 189

1250 1200 1125 1150

181 174 163 167

... ... ... ...

ST + AGED 480 °C (900 °F), 1 H ST + AGED 510 °C (950 °F), 1 H ST + AGED 550 °C (1025 °F), 1 H AGED 480 °C (900 °F), 1 H ST + AGED 480 °C (900 °F), 1 H AGED 540 °C (1000 °F), 1 H ST + AGED 540 °C (1000 °F), 1 H AGED 595 °C (1100 °F), 1 H ST + AGED 595 °C (1100 °F), 1 H AGED 480 °C, (900 °F), 1 H ST + AGED 480 °C (900 °F), 1 H AGED 540 °C (1000 °F), 1 H ST + AGED 540 °C (1000 °F), 1 H

1380 1280 1195 1225 1290 1105 1170 1095 1090 1090 1280 1070 1150

200 186 173 178 187 160 170 159 158 158 186 155 167

1275 1205 1115 1060 1185 940 1110 770 925 815 1125 745 1005

185 175 162 154 172 136 161 112 134 118 163 108 146

... ... ... ... ... ... ... ... ... ... ... ... ...

TREATMENT

IN.

REDUCTION IN AREA, %

FAILURE LOCATION(C)

IZOD IMPACT ENERGY(D) J FT · LBF

3.0 5.0 5.5 5.0 7.0 9.0 9.0 9.0 9,2

... ... ... ... ... ... 35 42 45

FL BM FL FL WM BM WM WM WM

... ... ... ... ... ... ... ... ...

... ... ... ... ... ... ... ... ...

14.0 14.0 15.0 ...

7.0 7.5 7.5 7.3

... ... ... 21

WM WM WM WM

... ... ... 15

... ... ... 11

13.0 12.0 12.0 10.0 11.0 11.0 12.0 13.5 12.0 ... ... ... ...

6.5 6.0 6.0 ... ... ... ... ... ... 8.0 9.5 9.5 9.5

29 30 32 32 48 29 48 50 54 21 36 27 39

... ... ... BM WM WM WM BM WM WM WM WM WM

... ... ... ... ... ... ... ... ... 38 8 46 23

... ... ... ... ... ... ... ... ... 28 6 34 17

GTA WELDS(A)

1.3 1.3 1.3 2.3 2.3 2.3 4.8 4.8 4.8

0.050 0.050 0.050 0.090 0.090 0.090 0.188 0.188 0.188

GMA WELDS(A)

4.8 4.8 4.8 25

0.188 0.188 0.188 1.00

SMA WELDS(A)

4.8 4.8 4.8 9.5 9.5 9.5 9.5 9.5 9.5 25 25 25 25

0.188 0.188 0.188 0.375 0.375 0.375 0.375 0.375 0.375 1.00 1.00 1.00 1.00

25 25

1.00 1.00

AGED 595 °C (1100 °F), 1 H ST + AGED 595 °C (1100 °F), 1 H

1090 1090

158 158

765 890

111 129

... ...

... ...

7.5 11.0

18 38

WM WM

Source: Ref 25

(A) (B) (C) (D)

ALL WELDS MADE IN CONDITION A WITH A 17-4PH COATED ELECTRODE OR FILLER WIRE. ST = SOLUTION TREATED AT 1035 °C (1900 °F) FOR 30 MIN; AIR COOLED TO 15 °C (60 °F) AND HELD FOR 1 H. WM, WELD METAL; FL, FUSION LINE; BM, BASE METAL. NOTCH IN TRANSVERSE FACE OF WELD METAL.

12 30

9 22

Base-Metal Condition for Welding. Depending on the thickness and level of restraint, 17-4PH stainless steel can be

welded in the annealed (A) or overaged (H1100 or H1150) conditions. The welding of material in conditions H900 through H1075 is generally not recommended. Material up to 25 mm (1.0 in.) thick can usually be welded in condition A provided that the weld is not very heavily restrained. Highly restrained welds or welds in heavier sections are best welded in conditions H1100 or H1150. Filler Metals. Type 630 covered electrodes and filler wires, which deposit weld metal that closely matches 17-4PH base

metal in composition and heat treatment response, are used when high-strength welds are required. Austenitic filler metals, such as type 308L, can be used for joining 17-4PH when high strength is not required in the joint. Welds made with austenitic filler metals are less susceptible to hydrogen embrittlement than those made with martensitic filler metals and may be used in the as-welded condition. Postweld Heat Treatment. In the as-welded condition, the weld metal and the high-temperature regions of the HAZs

of welds in 17-4PH stainless steel have structures consisting primarily of untempered martensite plus a small amount of ferrite (Fig. 2a). Weldments exhibit an aging peak in the HAZ and a weld metal hardness that is only slightly less than that of the base metal in condition A (Fig. 3a). Weldments in 17-4PH stainless steel are not usually put into service in the as-welded condition except for repair welds where PWHT is impractical. In order to obtain weld properties approximating those of the base metal, PWHT is necessary. For single-pass welds made with the base metal in condition A, a simple aging treatment of 1 to 4 h at 480 to 620 °C (900 to 1150 °F) is usually sufficient. It simultaneously hardens the weld metal, HAZ, and base metal and lowers the residual stresses associated with the weld. Because only slight overaging occurs in the portion of the HAZ that is heated into the aging temperature range during welding (Fig. 3a), joint efficiencies of 97 to 100% are obtained (Ref 25).

FIG. 3 MICROHARDNESS PROFILES OF AS-WELDED AND POSTWELD HEAT-TREATED WELDS IN 17-4PH MADE WITH (A) PAW AND (B) EBW PROCESSES. SOURCE: REF 26

In multipass welds, the repeated heating from the deposition of successive beads may leave a variation in structure from bead to bead that will result in nonuniform response to the aging treatment. Consequently, in the aged condition, weld

yield and ultimate tensile strengths are only about 65% and 80 to 90%, respectively, of the base-metal values (Ref 25). Solution treating the weld before hardening reduces the weld metal and HAZ ferrite contents and improves weld metal uniformity and response to heat treatment. As a result, weld strength increases to 80 to 90% and 90 to 95%, respectively, of the base-metal yield and ultimate tensile strengths. For welds made with the base metal in the overaged condition, solution treatment is required if it is desired to heat treat the weldment to a higher strength level. In general, if the weld deposit is less than 13 mm (0.5 in.) in thickness, fairly good tensile properties can be obtained even if the solution treatment is omitted prior to aging. However, the toughness of the weld metal decreases with aging temperature above 540 °C (1000 °F), probably due to an unfavorable carbide morphology. Therefore, if weld deposits are 13 mm (0.5 in.) or greater in thickness and a postweld solution treatment is not feasible, an age-hardening temperature of 550 °C (1025 °F) or lower is suggested (Ref 25). Electron-beam (EB) welding is used for joining 17-4PH stainless steel when precision welds with maximum strength and minimum distortion are required. Welding parameters are similar to those used for welding the standard austenitic stainless steels and the same considerations as in arc welding apply to the choice of base-metal condition for welding. The most important characteristic of an EB weld is the total heat input, which is only about 20% of that of a comparable arc weld (Ref 26). As a result, the HAZ is much narrower (Fig. 3b) and the weld metal is finer grained and has a lower ferrite content. Furthermore, since an EB weld is usually a single-pass weld, the weld-metal structure does not exhibit the inhomogeneities of a multipass arc weld. Consequently, the mechanical properties of EB welds are generally superior to those of arc welds, particularly in regard to fracture toughness (Ref 27). Typical mechanical properties of EB welds in 174PH are presented in Table 7. Postweld aging alone results in joint efficiencies of 95 to 100% (Ref 26, 28) and total elongation values only slightly lower than those of the base metal; however, the toughness of welds in the H900 condition is low (Ref 26). Re-solution treating before aging has little effect on weld strength but markedly increases the toughness of welds aged at 480 °C (900 °F). However, it produces only a small increase in toughness for welds aged at higher temperatures.

TABLE 7 MECHANICAL PROPERTIES OF ELECTRON-BEAM WELDS IN MARTENSITIC PH STEELS THICKNESS

PREWELD CONDITION

POSTWELD HEAT TREATMENT(A)

mm in.

TENSILE STRENGTH

0.2% YIELD STRENGTH

ELONGATION IN 50 MM (2 IN.), %

FAILURE LOCATION(B)

CVN IMPACT ENERGY j ft · lbf

FRACTURE TOUGHNESS mpa m

ksi

REF

mpa

ksi

mpa

ksi

1435 1400

208 203

1365 1270

198 184

11.0 12.2

BM BM

2.4 14.0

1.8 10.3

... ...

... ...

26 26

1165 1140

169 165

1125 1090

163 158

12.0 12.0

BM BM

15.6 18.3

11.5 13.5

... ...

... ...

26 26

1025 1020

149 148

985 985

143 143

15.4 14.6

BM BM

17.6 19.4

13.0 14.3

... ...

... ...

26 26

965

140

745

108

16.4

BM

20.6

15.2

...

...

26

1315 885 1055

191 128 153

1235 705 ...

179 102 ...

12.6(C) 11(D) TO 22 ...

... WM/BM ...

... ... ...

... ... ...

... ... 28 ... ... 27 KQ = KQ = 27 187 170 KQ = KQ = 27 194 176 KIC = KIC = 27 262 238

in

17-4 PH

3.2 3.2

0.125 A 0.125 A

3.2 3.2

0.125 A 0.125 A

3.2 3.2

0.125 A 0.125 A

3.2

0.125 A

9.5

0.375 A (E) H1150 (E) H1150

(E) (E)

AGED 480 °C (900 °F), 1 H ST + AGED 480 °C (900 °F), 1H AGED 550 °C (1025 °F), 4 H ST + AGED 550 °C (1025 °F), 4 H AGED 620 °C (1150 °F), 4 H ST + AGED 620 °C (1150 °F), 4 H AGED 760 °C (1400 °F), 2 H + 620 °C (1150 °F), 4 H AGED 480 °C (900 °F), 1 H NONE ST + AGED 595 °C (1 100 °F), 4 H

15-5PH (E)

(E)

1150

ST + AGED 565 °C (1050 °F), 4 H

...

...

1125

163

...

...

...

...

(E)

(E)

1150

ST + AGED 595 °C (1100 °F), 4 H

...

...

1020

148

...

...

...

...

H1150M

ST + AGED 540 °C (1000 °F), 4 H

...

...

1415

205

...

...

...

...

PH13-8 MO (E)

(A) (B) (C) (D) (E)

(E)

KIC = KIC = 31 110 100

ST = SOLUTION TREATED: 17-4PH AND 15-5PH AT 1035 °C (1900 °F), PH 13-8 MO AT 925 °C (1700 °F) FOR 30 MIN; AIR COOL. WM, WELD METAL; BM, BASE METAL. IN 25 MM (1 IN.). LOW VALUE IS FOR FAILURE IN THE WELD METAL. EXACT VALUE NOT REPORTED. IN RANGE OF 23 TO 60 MM (0.90 TO 2.37 IN.)

15-5PH Stainless Steel 15-5PH stainless steel differs from 17-4PH in that its composition is balanced so that its structure is ferrite-free and material for bar and forging billets is vacuum arc remelted. As a result, the transverse ductility and toughness of 15-5PH stainless steel are superior to those of 17-4PH. However, its weldability is essentially the same as that of 17-4PH and the same welding processes and procedures are used. A matching 15-5PH filler wire (AMS 5826) is available for critical applications but it is most often welded with type 630 (17-4PH) covered electrodes or filler wire. Postweld heat treatment and weld mechanical properties are essentially identical to those for 17-4PH. Autogenous welds in 15-5PH stainless steel contain less δ-ferrite than those in 17-4PH. Sufficient δ-ferrite is normally formed to prevent weld solidification cracking in GTA welds. However, the rapid cooling rates that are characteristic of EB and laser welding can produce welds with little or no δ-ferrite, which are susceptible to weld solidification cracking if welded with excessively high welding speeds (Ref 29). The strengths of defect-free EB welds in 15-5PH are similar to those in 17-4PH, but the fracture toughness is higher (Table 7). PH13-8 Mo Stainless Steel PH13-8 Mo stainless steel is vacuum induction melted plus vacuum arc remelted. It is ferrite-free and is used where high fracture toughness is required in heavy sections. Weldability is similar to 17-4PH stainless steel and the same procedures are used. It can be welded with the same processes, although covered electrodes for SMA welding are not available. Gastungsten arc and GMA welds can be made using a matching filler wire (AMS 5840). Although the final weld metal microstructure is ferrite-free (Fig. 2b), PH13-8 Mo has excellent resistance to weld solidification cracking because it solidifies as δ-ferrite and has low sulfur and phosphorus contents (Ref 30). After postweld solution treating and aging, the yield strengths and Charpy V-notch (CVN) impact energies of multipass GTA welds in PH13-8 Mo stainless steel are 96 to 100% and 50 to 100%, respectively, of the base-metal values (Fig. 4a), the difference between weld- and base-metal properties being greatest for aging temperatures above 540 °C (1000 °F). In heavy sections, PH13-8 Mo stainless steel is often joined by EB welding. After PWHT, EB welds have joint efficiencies of virtually 100% and their fracture toughness is approximately the same as that of the base metal (Ref 31).

FIG. 4 EFFECT OF AGING TEMPERATURE ON THE YIELD STRENGTH AND CHARPY V-NOTCH IMPACT ENERGY OF (A) PH13-8 MO AND (B) CUSTOM 450 BASE METAL AND WELD METAL. GTA WELDS MADE IN 13 MM (0.50 IN.) THICK PLATE (CONDITION A) WITH MATCHING FILLER METALS. CLOSED SYMBOLS REPRESENT YIELD STRENGTH VALUES; OPEN SYMBOLS REPRESENT IMPACT VALUES. SOURCE: REF 30

Custom 450 Stainless Steel The composition of Custom 450 stainless steel is balanced so that base metal, HAZ, and weld metal are essentially fully martensitic. In most respects, its weldability is similar to that of 17-4PH, and it can be joined using the same welding processes and procedures (Ref 32). A matching filler wire (AMS 5763) is usually used for GTA and GMA welding Custom 450 and is essential if the weldment is to be used without PWHT. A type 630 (17-4PH) filler wire or coated electrode can be used for joining Custom 450 if the weldment is postweld aged or solution treated and aged.

Custom 450 stainless steel is often used in condition A when only moderate strength is required. Material welded in this condition has acceptable weld properties in the as-welded condition, but CVN impact energy is substantially increased by a postweld solution anneal at 1040 °C (1900 °F) (Ref 32). Higher strength levels are obtained by postweld aging or solution treating and aging. After postweld solution treating and aging, the yield strengths of multipass GTA welds in Custom 450 are 98 to 100% of the base-metal values. Weld metal CVN impact energy increases with aging temperature from about 50% of the base-metal value for welds aged at 480 °C (900 °F) to 100% for welds aged at 565 °C (1050 °F) as shown in Fig. 4(b). Custom 455 Stainless Steel Custom 455 stainless steel is vacuum induction melted plus vacuum arc remelted and is completely martensitic. It is joined primarily by GTA welding and a Custom 455 filler metal is available for multipass welds. Procedures are similar to those used for 17-4PH, but a lower postweld solution treatment temperature (815 to 845 °C, or 1500 to 1550 °F) is used. In small sections, joint efficiencies of 80% are obtained by postweld aging only and 100% by solution treating and aging the weldment (Ref 33, 34). In large sections, joint efficiencies in excess of 85% are possible (Ref 19).


19. "CARPENTER CUSTOM 455," PRODUCT DATA BULLETIN, CARPENTER TECHNOLOGY CORP., READING, PA, 1971 21. B. POLLARD, UNPUBLISHED WORK 22. N. SUUTALA, EFFECT OF SOLIDIFICATION CONDITIONS ON THE SOLIDIFICATION MODE IN STAINLESS STEELS, METALL. TRANS. A, VOL 14, FEB 1983, P 191-197 23. G.E. LINNERT, WELDING PRECIPITATION-HARDENING STAINLESS STEELS, WELD. J., VOL 36, JAN 1957, P 9-27 24. F.J. HAWKINS, WELDING OF AGE-HARDENABLE STAINLESS STEELS, WELD. RES. COUNC. BULL., NO. 103, FEB 1965 25. "ARMCO FABRICATING DATA," BULLETIN SF-2, ARMCO, INC., MIDDLETOWN, OH, FEB 1966 26. C. CHEN AND J.S. YEH, ELECTRON BEAM AND PLASMA ARC WELDING OF 17-4PH STAINLESS STEEL, PROC. CONF. POWER BEAM PROCESSING: ELECTRON, LASER, PLASMA ARC (SAN DIEGO, CA), ASM INTERNATIONAL, 1988, P 45-54 27. A.J. TURNER, ELECTRON BEAM WELDING THICK SECTION PRECIPITATION-HARDENING STEEL, WELD. J., VOL 60, JAN 1981, P 18-26 28. M.U. ISLAM, G. CAMPBELL, AND R. HSU, FATIGUE AND TENSILE PROPERTIES OF EB WELDED 17-4PH STEEL, WELD. J., VOL 68, SEPT 1989, P 45-50 29. S. MATSUI ET AL., ELECTRON BEAM WELDING OF THE PRECIPITATION HARDENING STAINLESS STEEL, REVUE DE LA SOUDURE-LASTIJDSCRIFT, VOL 46, JAN 1990, P 52-57 30. W.R. CIESLAK, J.A. BROOKS, AND W.M. GARRISON, JR., THE WELDABILITY, MICROSTRUCTURE AND PROPERTIES OF PRECIPITATION STRENGTHENED MARTENSITIC STAINLESS STEELS, PROC. CONF. ADVANCES IN WELDING SCIENCE AND TECHNOLOGY (GATLINBURG, TN), ASM INTERNATIONAL, 1986 31. A.J. TURNER, ELECTRON BEAM WELDING OF HIGH VOLUME AIRCRAFT COMPONENTS, PROC. CONF. FUNDAMENTAL AND PRACTICAL APPROACHES TO THE RELIABILITY OF WELDED STRUCTURES (OSAKA, JAPAN), VOL 1, JAPAN WELDING SOCIETY, 1982, P 139-144 32. "WELDING CARPENTER CUSTOM 450 MULTIPLE-STRENGTH STAINLESS STEEL," PRODUCT DATA BULLETIN, CARPENTER TECHNOLOGY CORP., READING, PA, JAN 1973 33. ALLOY DIG., FILING CODE SS-193, APRIL 1967 34. R. KALTENHAUSER, WELDABILITY OF PRECIPITATION HARDENING STEELS, MET. ENG. QUART., VOL 9 (NO. 1), FEB 1969, P 44-57

Selection of Wrought Precipitation-Hardening Stainless Steels B. Pollard

Welding Semiaustenitic PH Steels As shown in Tables 1, 2, 3, and 4, semiaustenitic PH stainless steels include types 17-7PH, PH15-7 Mo, AM-350, and AM-355. As with the previous section on martensitic PH stainless steels, both microstructural evolution and weld parameters associated with semiaustenitic PH steels are described. The article "Welding of Stainless Steels" in this Volume should also be consulted. Microstructural Evolution The semiaustenitic PH steels have a primary δ-ferritic solidification mode and, on cooling, the δ-ferrite partially transforms to austenite by a Widmanstätten reaction. This austenite is stable down to ambient temperature so the aswelded structure of the fusion zone consists of austenite in a ferrite matrix (Fig. 2c). With the exception of AM-355, the percentage of retained ferrite in semiaustenitic PH welds is substantially higher than for welds in martensitic PH steels. Postweld conditioning treatments destabilize the austenite so that it transforms to martensite at ambient or subzero temperatures. Subsequent aging simultaneously tempers the martensite and precipitates the strengthening particles, which are not visible by optical microscopy. The final fusion zone microstructure, therefore, consists of areas of retained ferrite in a tempered martensitic matrix. If a low conditioning temperature is used, carbide particles may also be observed at the ferrite-martensite interfaces. In the as-welded and, to a lesser degree, the postweld heat-treated conditions, the HAZ is distinguishable from the base metal by an increase in grain size and ferrite content in the high-temperature region close to the fusion line. 17-7PH and PH15-7 Mo Stainless Steels Weldability. 17-7PH and PH15-7 Mo stainless steels are typical of the semiaustenitic PH steels. As-deposited welds

made without filler metal have weld metal and HAZ structures that are predominantly austenitic (Fig. 2c). This ductile structure is not susceptible to cold cracking, so 17-7PH and PH15-7 Mo are readily welded in any condition without preheat and without the need for critical control of interpass temperature or cooling rate (Ref 24, 35). However, autogenous 17-7PH weld metal is not fully austenitic, but rather contains about 25% δ-ferrite versus 10 to 15% in the base metal. Likewise, autogenous PH15-7 Mo weld metal may contain up to 40% δ-ferrite versus 15% in the base metal. The δ-ferrite content of the weld metal is determined primarily by composition, but high values are also promoted to a lesser degree by thicker weld beads, higher welding currents, and higher welding speeds. A large volume fraction of δ-ferrite can make the weld metal susceptible to hot cracking and reduce the ductility of the weld metal in the heat-treated condition. If necessary, the δ-ferrite content of the weld metal can be reduced by the addition of an appropriate filler metal. Filler Metals. Filler wires are available for 17-7PH (AMS 5824) and PH15-7 Mo (AMS 5812) stainless steels. The

compositions of these wires have been adjusted so as to produce deposits with lower -ferrite contents than would be produced by the base-metal compositions while still responding correctly to the age-hardening treatment (Ref 24, 35). Covered electrodes are not available for 17-7PH and PH15-7 Mo stainless steels because excessive oxidation of aluminum would occur in the arc. However, type 630 (17-4PH) covered electrodes have been used successfully for joining 17-7PH. To obtain adequate heat treat response with this filler metal, welding conditions that produce high dilution and a compromise PWHT practice are required (Ref 24). A heat treatment of 870 °C (1600 °F) for 90 min is used to simultaneously solution treat the type 630 weld metal and condition the 17-7PH base metal (Ref 35). The weld is then refrigerated at -75 °C (-100 °F) for 8 h and aged at 510 to 595 °C (950 to 1050 °F) for 1 h, depending on the properties required. Standard iron-chromium-nickel filler metals, such as types 308, 308L, 309, and 347, can also be used for joining 17-7PH and PH15-7 Mo, but the weld metal will not respond to heat treatment in the same manner as the base metal. Response to heat treatment will depend on the ratio of austenite-to-ferrite-forming elements in the filler metal and the amount of dilution by the base metal. In most instances, weld strength will be much lower than that obtained with a matching filler metal, but ductility and toughness will be higher. Gas-Tungsten Arc Welding. Because 17-7PH and PH15-7 Mo stainless steels are used primarily as sheet and strip,

they are most frequently joined by GTA welding. Preheating and postheating are not required to prevent cracking, but the

selection of welding conditions requires more care than for the standard austenitic stainless steels. Helium gas is preferred over argon for shielding the arc because helium minimizes the formation of a continuous aluminum oxide film on the surface of the weld pool (Ref 23). Particular care should be taken to provide efficient shielding of the weld pool in order to minimize loss of aluminum. Inert gas shielding should also be used on the underside of full penetration welds to protect the weld root. Welding speeds typically are of the order of 5 mm/s (12 in./min) (Ref 35). Higher welding speeds can cause porosity at the fusion line. Welding currents and weld-pass sequence should be selected to minimize the weld crosssectional area. Gas-metal arc welding is used for joining 17-7PH and PH15-7 Mo in heavy sheet and plate thicknesses. A mixture of 75He-25Ar is the most widely used shielding gas. Welding conditions are otherwise similar to those used for the standard austenitic stainless steels. Postweld Heat Treatment. As-deposited welds in 17-7PH and PH15-7 Mo steels have a structure of austenite plus

ferrite that is of much lower strength than the hardened base metal. A PWHT is, therefore, essential if a weld strength similar to the base-metal strength is required. The effects of PWHT on the tensile properties of GTA welds in 17-7PH and PH15-7 Mo steels are presented in Table 8; corresponding welding conditions are given in Table 9. The joint efficiencies of welds postweld hardened to the RH950 or TH1050 conditions without reannealing are 100% with weld reinforcement and 95 to 98% without it. Postweld reannealing before hardening has negligible effect on weld strength but increases weld ductility.

TABLE 8 TENSILE PROPERTIES OF GTA WELDS IN 1.6 MM (0.062 IN.) THICK 17-7PH AND PH15-7 MO STEELS See also Table 9 ALLOY

PREWELD CONDITION(A)

POSTWELD HEAT TREATMENT(A)

HARDENED AND TESTED AS-WELDED PH15-7 MO TH1050 NONE PH15-7 MO R950 NONE WELDED AND HARDENED 17-7PH A TH1050 17-7PH A RH950 PH15-7 MO A TH1050 PH15-7 MO A TH1050 PH15-7 MO A RH950 PH15-7 MO A RH950 WELDED, REANNEALED, AND HARDENED 17-7PH A A + TH1050 PH15-7 MO A A + TH1050 PH15-7 MO A A + TH1050 PH15-7 MO A A + RH950 PH15-7 MO A A + RH950 PH15-7 MO A R + A + RH950

REINFORCEMENT

MPA

KSI

0.2% YIELD STRENGTH MPA KSI

OFF OFF

840 795

122 115

605 600

ON ON ON OFF ON OFF

1415 1515 1450 1415 1655 1585

205 220 210 205 240 230

ON ON OFF ON OFF OFF

1435 1450 1380 1620 1585 1660

208 210 200 235 230 241

TENSILE STRENGTH

ELONGATION, %

FAILURE LOCATION(B)

IN 13 MM (0.5 IN.)

IN 50 MM (2 IN.)

88 87

25.0 22.0

7.0 6.0

WM WM

1345 1415 1415 1345 1550 1480

195 205 205 195 225 215

6.0 6.0 ... 8.0 ... 5.0

1.5 2.0 5.0 2.0 5.0 2.0

BM BM BM WM BM WM

1345 1415 1345 1480 1450 1505

195 205 195 215 210 218

8.0 ... 10.0 ... 7.0 9.0

2.5 5.0 2.5 4.0 2.0 3.5

BM BM WM BM WM WM

(A) PREWELD AND POSTWELD HEAT TREATING CONDITIONS. A (CONDITION A): ANNEAL AT 1065 °C (1950 °F) FOR 10 MIN. AIR COOL. R (REFRIGERATION TREATMENT FOR TRANSFORMATION): -75 °C (-100 °F) FOR 8 H. TH1050: 760 °C (1400 °F) FOR 90 MIN, AIR COOL + 565 °C (1050 °F) FOR 90 MIN, AIR COOL. RH950: 955 °C (1750 °F) FOR 10 MIN; AIR COOL + -75 °C (-100 °F) FOR 8 H + 510 °C (950 °F) FOR 60 MIN; AIR COOL. (B) WM, WELD METAL; BM, BASE METAL

TABLE 9 WELDING CONDITIONS FOR GTA WELDS DESCRIBED IN TABLE 8

WELDING CONDITION THICKNESS, MM (IN.) SHIELDING GAS GAS FLOW RATE, L/RAIN (FT3/H) WELDING SPEED, MM/S (IN./MIN) CURRENT FILLER WIRE WIRE FEED RATE, MM/S (IN./MIN)

ALLOY 17-7PH 1.59 (0.0625) HELIUM 28 (60)

PH15-7 MO 1.57 (0.062) HELIUM 24 (50)

6.3 (15)

5.1 (12)

AS REQUIRED FOR FULL PENETRATION IN A SINGLE PASS 1.6 MM (0.062 IN.) DIAM W17-7PH

AS REQUIRED FOR FULL PENETRATION IN A SINGLE PASS 1.6 MM (0.062 IN.) DIAM WPH157MO 1 TO 1

12 (28)

Resistance Welding. Spot and seam welding have been used extensively for joining 17-7PH and PH15-7 Mo steels. For good results, particular attention should be given to the surface cleanliness of the steel. Surfaces to be joined should be free of oil, grease, dirt, and heat treating scale. Welding schedules should use longer weld times, lower weld currents, and electrode forces 10 to 20% higher than those normally used for the standard austenitic stainless steels (Ref 24). The best mechanical properties are obtained by welding just before or after final hardening. This practice produces weld nuggets with a tough austenitic microstructure. The tensile strength normal to the weld interface is a maximum for this condition, and the shear strength is not significantly lower than for nuggets in the hardened condition (Ref 35).

Sound flash butt welds can be produced in 17-7PH and PH15-7 Mo stainless steels, but ductility may be low as a result of ferrite stringers in the base metal being bent parallel to the weld interface (Ref 35). AM-350 and AM-355 Stainless Steels Weldability. AM-350 and AM-355 differ from 17-7PH and PH15-7 Mo stainless steels in that they do not contain

aluminum. Martensitic transformation provides their major strengthening mechanism with an additional but smaller contribution from the precipitation of Cr2N or an isostructurally similar phase during tempering (Ref 12). The absence of any reactive elements makes the weldability of AM-350 and AM-355 as good as that of the standard austenitic stainless steels and superior to that of 17-7PH. Although these steels are most often joined by GTA and GMA welding, they can be joined by any of the processes used for the standard austenitic stainless steels, including submerged arc, electron-beam, laser-beam, and resistance welding (Ref 34, 36, 37). Like 17-7PH and PH15-7 Mo, these steels may be welded in the annealed or hardened conditions and neither preheat nor special precautions with regard to interpass temperature or cooling rate are necessary. However, PWHT must be used if maximum weld strength is required. Annealed AM-350 base metal contains 10 to 15% δ-ferrite, and welding produces even higher levels in autogeneous weld metal and the high-temperature portion of the HAZ. In the subzero cooled and tempered (SCT) condition, the final structure consists of islands of ferrite surrounded by carbides in a tempered martensitic matrix. In contrast, the composition of AM-355 is balanced so that the base metal contains less than 5% δ-ferrite in condition H (solution treated at 1065 °C, or 1950 °F) and 0% ferrite in the equalized plus overaged and tempered (EOT+) condition, and welding does not increase the ferrite content to any significant extent. Filler Metals. An AM-355 filler wire (AMS 5780) with a composition identical to that of the base metal is available for

joining AM-355. Likewise, a matching AM-350 filler wire (AMS 4774) is available for welding AM-350, but the AM355 filler wire is usually preferred for this material because it gives slightly higher weld strengths after PWHT (Ref 36). The AM-355 filler wire can also be used with a proprietary flux for submerged arc welding. AM-355 covered electrodes have been developed for SMA welding but may not be readily available. Postweld Heat Treatment. Joint efficiencies of 90 to 100% can be obtained by PWHT of AM-350 and AM-355 welds

(Ref 36). The PWHT for welds in AM-350 consists of annealing at 930 °C (1710 °F) (condition L) and cooling rapidly to

room temperature, followed by subzero cooling to -75 °C (-100 °F) and tempering at 455 or 540 °C (850 or 1000 °F) (condition SCT). Unlike AM-350, AM-355 must have a suitably prepared microstructure prior to the condition L treatment in order to minimize intergranular precipitation during this treatment. The required microstructure is produced by a prior heat treatment which involves carbide solution annealing for 1 to 3 h at 1040 °C (1900 °F), cooling rapidly enough to room temperature to avoid intergranular carbide precipitation in the temperature range 870 to 650 °C (1600 to 1200 °F), and subzero cooling to -75 °C (-100 °F). The subzero cooling causes a partial transformation to martensite and, during subsequent heat treatments in the carbide precipitation range, carbides precipitate in the former martensitic areas in preference to the grain boundaries. After this treatment, AM-355 weldments are annealed at 915 to 970 °C (1675 to 1775 °F), subzero cooled to -75 °C (-100 °F), and tempered at 455 or 540 °C (850 or 1000 °F) (condition SCT). Resistance Welding. AM-350 and AM-355 can be spot welded either before or after the hardening treatment using

similar conditions as for standard austenitic stainless steels. Postweld heat treatments (for example, condition SCT) produce a moderate increase in tensile shear strength, but also a corresponding decrease in cross-tension strength. Consequently, the best combination of tensile shear and cross-tension strength is obtained by welding in the hardened condition and not postweld heat treating (Ref 24, 37).


12. G. AGGEN, "PHASE TRANSFORMATIONS AND HEAT TREATMENT STUDIES OF A CONTROLLED TRANSFORMATION STAINLESS STEEL ALLOY," DR. ENG. SC. THESIS, RENSSELAER POLYTECHNIC INSTITUTE, AUG 1963 23. G.E. LINNERT, WELDING PRECIPITATION-HARDENING STAINLESS STEELS, WELD. J., VOL 36, JAN 1957, P 9-27 24. F.J. HAWKINS, WELDING OF AGE-HARDENABLE STAINLESS STEELS, WELD. RES. COUNC. BULL., NO. 103, FEB 1965 34. R. KALTENHAUSER, WELDABILITY OF PRECIPITATION HARDENING STEELS, MET. ENG. QUART., VOL 9 (NO. 1), FEB 1969, P 44-57 35. "ARMCO FABRICATING DATA," BULLETIN SF-5, ARMCO, INC., MIDDLETOWN, OH, DEC 1967 36. "AM-350/AM-355 PRECIPITATION HARDENING STAINLESS STEELS," PRODUCT DATA BULLETIN, ALLEGHENY LUDLUM CORP., PITTSBURGH, PA, 1963 37. H.W. MISHLER, R.E. MONROE, AND P.J. RIEPPEL, "WELDING OF HIGH STRENGTH STEELS FOR AIRCRAFT AND MISSILE APPLICATIONS," DMIC REPORT 118, DEFENSE METALS INFORMATION CENTER, COLUMBUS, OH, 12 OCT 1959 Selection of Wrought Precipitation-Hardening Stainless Steels B. Pollard

Welding Austenitic PH Steels Weldability and welding recommendations for A-286 and JBK-75 austenitic PH stainless steels are discussed below. Tables 1, 2, 3, and 4 list properties and heat treatments for these steels. As noted previously, additional information can be found in the article "Welding of Stainless Steels" in this Volume. A-286 Stainless Steel Weldability. A-286 stainless steel is highly alloyed and since welds solidify entirely as austenite (Fig. 2d), they are

susceptible to hot cracking in both the fusion zone and HAZ (Ref 38, 39, 40, 41, 42). Solidification cracking in the fusion zone results from the formation of a low-melting-point Laves-phase eutectic at the grain boundaries, which is produced by the segregation of titanium and other solutes during the last stages of solidification (Ref 40, 42). This phase is enriched

in titanium, nickel, silicon, and molybdenum and depleted in iron and chromium (Ref 40). The same phase is also responsible for HAZ cracking (Ref 39, 40, 41, 42). However, in the HAZ, the Laves-phase eutectic is formed by the constitutional liquation of (Ti,Mo)(C,N) inclusions (Ref 41, 42). Constitutional liquation is a result of the rapid, nonequilibrium heating conditions typical of arc welding, which cause the carbonitrides to decompose more rapidly than their constituent elements can diffuse into the matrix. The result is the formation of solute-rich zones of eutectic composition (and hence low melting point) around the decomposing carbonitrides. This liquid then wets the grain boundaries, and the grains are pulled apart by shrinkage stresses as the weld cools. Of the residual elements normally present in A-286, silicon and boron have been found to have the greatest effect on hot cracking in both the fusion zone and HAZ (Ref 40, 41, 42). The ease with which crack-free arc welds can be produced in A-286 depends on the grade of material, weld-section thickness, amount of restraint, the welding process used, the welding procedure, and the choice of filler metal. Both regular and welding grades of A-286 are produced. The welding grades (AMS 5858 and AMS 5895) have lower limits on manganese, silicon, sulfur, and phosphorus than the regular grades and exhibit greater resistance to hot cracking. Whenever possible, they should be specified in preference to the regular grades for applications that involve welding. Although A-286 is classified as a difficult-to-weld material, it can be joined successfully by GTA, EB, and resistance welding if the correct procedures are used. Gas-Tungsten Arc Welding. When GTA welding A-286, the best results are obtained with the base metal in the solution treated condition, low-heat inputs, and minimum restraint (Ref 24, 34). Autogenous crack-free GTA welds can be produced in thin sheet, but welds in thick sheet and plate require filler metal. A stringer bead technique should be used for multipass welds in heavy sections.

The choice of filler metal depends on the application. The regular A-286 filler wires (AMS 5804) have lower limits on sulfur and phosphorus contents but are otherwise similar in composition to the base metal. Deposits made with these wires are the most prone to hot cracking, so they have been used primarily for single-pass welds (Ref 34). Premium quality A-286 filler wires (AMS 5805) with lower carbon, manganese, silicon, phosphorus, and boron contents and limits on oxygen, nitrogen, and hydrogen levels are also available. These wires are produced from vacuum induction melted material by wire rolling instead of drawing (Ref 43, 44). They produce high-purity deposits that are less susceptible to hot cracking than those made using the regular filler wires and are used for critical applications such as jet engine repairs. Dissimilar metal filler wires are the usual choice for multipass welds if a matching filler metal is not essential because they exhibit greater resistance to hot cracking than A-286 weld metal. Hastelloy W has been used extensively for joining A-286 (Ref 24, 43). It has excellent resistance to hot cracking, but its stress-rupture strength is substantially lower than that of A-286 weld metal and it is not suitable for many high-temperature environments. Type 310 filler metal has been used for joining A-286 plate (Ref 24). It has very good corrosion resistance in many media but is nonhardening. AMS 5675 (AWS A5.14 Class ERNiCrFe-6) filler metal is significantly less crack sensitive than A-286 and is age hardenable, but it does not quite match the strength of A-286. Postweld Heat Treatment. Typical tensile properties of welds in A-286 are presented in Table 10. A-286 weld metal

does not age harden as rapidly or to the same extent as the base metal. A PWHT of 725 °C (1335 °F) for 48 h is required to increase the yield and ultimate tensile strengths to 93% and 89%, respectively, of the base-metal values (Ref 45). However, aging is much more rapid if the weld is first solution treated (Ref 46). For example, solution treating at 1080 °C (1975 °F) for 2 h and aging at 725 °C (1335 °F) for 16 h increases the yield and ultimate tensile strengths to 94% and 97%, respectively, of the base-metal values. If a high-temperature solution treatment is not feasible, weld strength can be increased by using a low-temperature multistep heat treatment consisting of 725 °C (1335 °F), 4 h plus 650 °C (1202 °F), 5 h plus 725 °C (1335 °F), 44 h (Ref 45). This heat treatment raises the yield and ultimate tensile strengths to 99% and 93%, respectively, of the base-metal values.

TABLE 10 TENSILE PROPERTIES OF WELDS IN AUSTENITIC PH STEELS THICKNESS

mm

FILLER METAL

PREWELD CONDITION(A)

SPECIMEN(B)

in.

TENSILE STRENGTH

0.2% YIELD STRENGTH

ELONGATION, %


COMMENTS

REF

mpa

ksi

mpa

ksi

in 25 mm (1 in.)

AGED 725 °C (1335 °F), 48 H AGED 725 °C (1335 °F), 48 H AGED 725 °C (1335 °F), 4 H + 650 °C (1200 °F), 5 H + 725 °C 1335 °F, 44 H AGED 725 °C (1335 °F), 4 H + 650 °C (1200 °F), 5 H + 725 °C (1335 °F), 44 ST 980 °C (1795 °F), 2 H + AGED 725 °C (1335 °F), 16 H NONE ST 980 °C (1795 °F), 2 H + AGED 725 °C (1335 °F), 16 H

952

138

710

103 16

...

29

...

45

850

123

659

96

...

35

NO CRACKING

45

1020

148

777

113 20

...

28

...

45

952

138

769

112 . . .

...

45

NO CRACKING

45

1069

155

708

103 30

...

43

...

46

724 1033

105 150

611 665

89 96

... ...

34 48

... 46 SCATTERED 46 MICROFISSURES IN WELD METAL

AGED 730 °C (1350 °F), 16 H AGED 730 °C (1350 °F), 16 H

1047

152

793

115 . . .

23.0

...

...

919

133

775

112 . . .

5.6

...

REAGED 730 °C (1350 °F),

889

130

...

...

9.4

...

MICROFISSURES 47 IN HAZ MICROFISSURES 47 IN HAZ

POSTWELD HEAT TREATMENT

in 50 mm (2 in.)

A-286 GTA WELDS

3.0

0.118 NONE ST

BM

3.0

0.118 NONE ST

W

3.0

0.118 NONE ST

BM

3.0

0.118 NONE ST

W

35

1.38

A-286

ST

BM

35 35

1.38 1.38

A-286 A-286

ST ST

W W

8

15 31

A-286 EB WELDS

3.2

0.125 NONE ST

BM

3.2

0.125 NONE ST

W

3.2

0.125 NONE ST + AGED

W

...

47

16 H JBK-75 GTA WELDS

3.2

0.125 NONE ST

BM

3.2

0.125 NONE ST

W

3.2

0.125 NONE ST

BM

3.2

0.125 NONE ST

W

3.2

0.125 NONE ST

BM

3.2

0.125 NONE ST

W

(A) (B) (C)

AGED 725 °C (1335 °F), 24 H AGED 725 °C (1335 °F), 24 H AGED 725 °C (1335 °F), 24 H + 650 °C (1200 °F), 24 H AGED 725 °C (1335 °F), 24 H + 650 °C ( 1200 °F), 24 H ST 950 °C (1740 °F), 24 H + 725 °C (1335 °F), 24 H ST 950 °C (1335 °F), 24 H + 650 °C (1200 °F), 24 H

L130(C) 164(C) . . .

...

...

...

...

...

52

800(C)

116(C) . . .

...

...

...

...

NO CRACKING

52

1215

176

965

140 17

...

41

...

52

945

137

841

122 4

...

26

NO CRACKING

52

980(C)

142(C) . . .

...

...

...

...

...

52

980(C)

142(C) . . .

...

...

...

...

NO CRACKING

52

ST = SOLUTION TREATED. TRANSVERSE TENSILE SPECIMENS: BM. BASE METAL; W, WELDMENT. ESTIMATED FROM HARDNESS

Electron-Beam Welding. Although EB welding is a low-heat input process compared to arc welding, the use of EB

welding does not guarantee freedom from hot cracking. However, crack-free EB welds can be made in A-286 if heat input, welding speed, and restraint are low. A-286 can be EB welded in both the solution treated and hardened conditions. However, resistance to hot cracking is greater in the solution treated condition. EB welds in A-286 are usually made without filler metal. In the as-welded condition, weld strength is about the same as for GTA welds in the same condition, but postweld aging is more effective in increasing weld strength than for GTA welds in A-286. For example, the yield and ultimate tensile strengths of EB welds made in 3.2 mm (0.128 in.) thick A-286 in the solution treated condition has yield and ultimate tensile strengths of 98% and 88%, respectively, of the base-metal values after aging at 720 °C (1330 °F) for 16 h (Ref 47). Resistance Welding. A-286 stainless steel can be joined by resistance spot and seam welding in both the soft and

hardened conditions, but the welds, as with those made by other fusion welding processes, are susceptible to hot cracking. However, some hot cracking can be tolerated in spot and seam welds since cracking of up to 5 vol% of the weld nugget has little effect on static or cyclic weld strength if it is concentrated at the center of the weld nugget (Ref 48). Hot cracking can be reduced or eliminated by using electrode forces significantly higher than those used for the standard austenitic stainless steels (Ref 24, 48). Welding conditions that slow down the cooling rate (for example, downsloping the weld current) have also been recommended as a means of eliminating cracking (Ref 48). Limited data suggest that longer weld times and lower weld currents than those used for standard austenitic stainless steels may also be beneficial. Spot welds made in hardened material without PWHT have the best combination of shear and cross-tension strengths (Ref 37). Seam welds in A-286 are also susceptible to hot cracking. Limited data indicate that the best results are obtained with a lower wheel speed, lower weld current and longer (on plus off) time than are generally used for the standard austenitic stainless steels. Increased electrode force and continuous (rather than interrupted) weld current may also be used to prevent cracking (Ref 49). Little information is available on welding conditions for flash butt welding A-286 because these conditions vary from machine to machine and application to application. However, A-286 is routinely joined by flash butt welding with excellent results. After postweld aging at 720 °C (1325 °F), weld properties are virtually equal to those of the base metal (Ref 50). JBK-75 Stainless Steel JBK-75 is a modification of A-286 with the low carbon, manganese, silicon, and boron contents of welding grade A-286 plus a higher nickel content. The beneficial effect of the higher nickel content is explained by the concept of shrinkage brittleness (Ref 51). According to this concept, during the last stages of solidification, a minimum amount of eutectic liquid is required to fill all the dendritic fissures formed by the eutectic films. If a critical, subminimum amount of eutectic forms, severe cracking occurs. On the other hand, if only a very small amount of eutectic forms, no cracking occurs because enough solid boundaries are present to withstand the shrinkage stresses. Therefore, the beneficial effect of the high nickel content of JBK-75 is due to it increasing the amount of eutectic formed so that any microfissures which nucleate are self-healing. Consequently, JBK-75 has greater resistance to hot cracking than A-286 and can be GTA welded in heavier sections without cracking. Mechanical properties of GTA welds in JBK-75 are presented in Table 10. Hardness measurements indicate that the weld ultimate tensile strength (UTS) is only about 70% of the base-metal value after postweld aging at 725 °C (1335 °F) for 24 h (Ref 52). Higher weld- and base-metal strengths and a joint efficiency, based on the UTS, of about 78% are obtained by a double aging treatment consisting of 725 °C (1335 °F), 24 h plus 650 °C (1200 °F), for 24 h. Homogenizing at 950 °C (1740 °F) before aging also increases weld-metal strength but simultaneously reduces the strength of the base metal. For example, a heat treatment consisting of 950 °C (1740 °F) for 24 h plus 725 °C (1335 °F) for 24 h produces a hardness of 30 HRC, which is equivalent to a UTS of about 965 MPa (140 ksi) in both the weld and base metal.


24. F.J. HAWKINS, WELDING OF AGE-HARDENABLE STAINLESS STEELS, WELD. RES. COUNC. BULL., NO. 103, FEB 1965

34. R. KALTENHAUSER, WELDABILITY OF PRECIPITATION HARDENING STEELS, MET. ENG. QUART., VOL 9 (NO. 1), FEB 1969, P 44-57 37. H.W. MISHLER, R.E. MONROE, AND P.J. RIEPPEL, "WELDING OF HIGH STRENGTH STEELS FOR AIRCRAFT AND MISSILE APPLICATIONS," DMIC REPORT 118, DEFENSE METALS INFORMATION CENTER, COLUMBUS, OH, 12 OCT 1959 38. J.J. VAGI AND D.C. MARTIN, WELDING OF HIGH-STRENGTH STAINLESS STEELS FOR ELEVATED TEMPERATURE USE, WELD. J., VOL 35, MARCH 1956, P 137S-144S 39. B.S. BLUM AND R.H. WITT, HEAT AFFECTED ZONE CRACKING IN A-286 WELDMENTS, WELD. J., VOL 42, AUG 1963, P 365S-370S 40. J.A. BROOKS AND R.W. KRENZER, PROGRESS TOWARD A MORE WELDABLE A-286, WELD. J., VOL 53, JUNE 1974, P 242S-245S 41. J.A. BROOKS, EFFECT OF ALLOY MODIFICATIONS ON HAZ CRACKING OF A-286 STAINLESS STEEL, WELD. J., VOL 53, NOV 1974, P 517S-523S 42. J.A. BROOKS, A.W. THOMPSON, AND J.C. WILLIAMS, WELD CRACKING OF AUSTENITIC STAINLESS STEELS--EFFECTS OF IMPURITIES AND MINOR ELEMENTS, PROC. CONF. PHYS. METALLURGY OF METAL JOINING (ST. LOUIS, MO), TMS/AIME, 1980, P 117-136 43. W.H. SIMON, S. BANERJEE, AND R.S. CREMISIO, METALLURGICALLY CONTROLLED WELDING FILLER WIRE--TRACE ELEMENT CONTROLLED TO PRODUCE CONSISTENTLY SOUND WELDS, PROC. INT. CONF.: THE EFFECTS OF RESIDUAL IMPURITY AND MICROALLOYING ELEMENTS ON WELDABILITY AND WELD PROPERTIES (LONDON, ENGLAND), WELDING INSTITUTE, PAPER 14, 1983 44. D.W. HARVEY, METALLURGICALLY CONTROLLED WELDING FILLER WIRE, WELD. REV., MAY 1988, P 80, 82, 122, 123 45. L.T. SUMMERS, "EFFECT OF CHEMICAL SEGREGATION ON THE MICROSTRUCTURE AND PROPERTIES OF A-286 WELDMENTS," PH.D. DISSERTATION, UNIVERSITY OF CALIFORNIA, BERKELEY, 1984 46. A. INOUE ET AL., FATIGUE STRENGTH OF ALLOY A-286 AT CRYOGENIC TEMPERATURE, PROC. CONF. FATIGUE '87, VOL 1 (CHARLOTTESVILLE, VA), ENGINEERING MATERIALS ADVISORY SERVICES, LTD., 1987, P 547-555 47. R.E. ROTH AND N.F. BRATKOVICH, CHARACTERISTICS AND STRENGTH DATA OF ELECTRON BEAM WELDS IN FOUR REPRESENTATIVE MATERIALS, WELD. J., VOL 41, MAY 1962, P 229S-240S 48. I.F. SQUIRES, DETERMINATION OF RESISTANCE SPOT WELDING CONDITIONS FOR SOME HEAT-RESISTING STAINLESS STEELS, BR. WELD. J., VOL 9, MARCH 1962, P 149-157 49. P.M. KNOWLSON, DETERMINATION OF RESISTANCE SEAM WELDING CONDITIONS FOR SOME HEAT RESISTING STEELS, BR. WELD. J., VOL 9, 1962, P 436-445 50. J.J. VAGI, R.M. EVANS, AND D.C. MARTIN, "WELDING PRECIPITATION-HARDENING STAINLESS STEELS," NASA TECHNICAL MEMORANDUM X-53582, 28 FEB 1967 51. B.I. MEDOVAR, ON THE NATURE OF WELD HOT CRACKING, TRANSLATED FROM AVTOM. SVARKA, VOL 7 (NO. 4), 1954, P 12-28 52. M.J. STRUM, L.T. SUMMERS, AND J.W. MORRIS, JR., THE AGING RESPONSE OF A WELDED IRON-BASED SUPERALLOY, WELD. J., VOL 62, SEPT 1983, P 235S-242S Selection of Wrought Precipitation-Hardening Stainless Steels B. Pollard

References

1. V. SEETHARAMAN ET AL., PRECIPITATION HARDENING IN A PH13-8 MO STAINLESS STEEL, MATER. SCI. ENG., VOL 47, JAN 1981, P 1-11 2. H.J. RACK AND D. KALISH, THE STRENGTH, FRACTURE TOUGHNESS AND LOW CYCLE FATIGUE BEHAVIOR OF 17-4PH STAINLESS STEEL, METALL. TRANS. A, VOL 5, JULY 1974, P 1595-1605 3. K.C. ANTHONY, AGING REACTIONS IN PRECIPITATION HARDENABLE STAINLESS STEEL, J. MET., VOL 15, DEC 1963, P 922-927 4. M. HENTHORNE, T.A. DEBOLD, AND R.J. YINGER, "CUSTOM 450-A NEW HIGHER STRENGTH STAINLESS STEEL," PAPER NO. 53, CORROSION '72, NATIONAL ASSOCIATION OF CORROSION ENGINEERS CONFERENCE, 1972 5. J.E. MCBRIDE, JR., AND G.N. MANIAR, METALLOGRAPHY AS A QUALITY CONTROL TOOL, PLENUM PUBLISHING, 1980, P 279 6. S. WIDGE, "A STUDY OF THE EFFECTS OF VARYING SOLUTION TREATMENT ON THE TOUGHNESS OF A STAINLESS MARAGING ALLOY," PH.D. DISSERTATION, LEHIGH UNIVERSITY, 1984 7. E.E. UNDERWOOD, A.E. AUSTIN, AND G.K. MANNING, THE MECHANISM OF HARDENING IN 17-7 NI-CR PRECIPITATION-HARDENING STAINLESS STEEL, JISI, VOL 200, AUG 1962, P 644651 8. H.C. BURNETT, R.H. DUFF, AND H.C. VACHER, IDENTIFICATION OF METALLURGICAL REACTIONS AND THEIR EFFECT ON MECHANICAL PROPERTIES OF 17-7PH STEEL, J. RES. NATL. BUR. STAND.--ENG. INSTRUMENTATION, VOL 66C, 1962, P 113-119 9. E.G. FEL'DGANDLER AND M.V. PRIDANTSEV, PHASE TRANSFORMATION IN KH 17 N7 YU STAINLESS STEEL, METALLOVED. TERM. OBRAB. MET., NOV 1960, P 2-7 10. J.C. WILKENS AND R.E. PENCE, A STUDY OF THE MICROSTRUCTURE OF PRECIPITATION HARDENING STAINLESS STEEL SHEET, ADVANCES IN ELECTRON METALLOGRAPHY AND ELECTRON PROBE MICROANALYSIS, STP 317, ASTM, 1962, P 140-149 11. J.C. WILKENS, PRIVATE COMMUNICATION, IN PRECIPITATION FROM IRON-BASE ALLOYS, PROCEEDINGS OF AN AIME SYMPOSIUM, CLEVELAND, OH, 21 OCT 1963, P 98 12. G. AGGEN, "PHASE TRANSFORMATIONS AND HEAT TREATMENT STUDIES OF A CONTROLLED TRANSFORMATION STAINLESS STEEL ALLOY," DR. ENG. SC. THESIS, RENSSELAER POLYTECHNIC INSTITUTE, AUG 1963 13. D.R. MUZYKA, THE METALLURGY OF NICKEL-IRON ALLOYS, THE SUPERALLOYS, C.T. SIMS AND W.C. HAGEL, ED., JOHN WILEY & SONS, 1972, P 113-143 14. A.W. THOMPSON AND J.A. BROOKS, THE MECHANISM OF PRECIPITATION STRENGTHENING IN AN IRON-BASE SUPERALLOY, ACTA METALL., VOL 30, DEC 1982, P 2197-2203 15. T.J. HEADLEY, M.M. KAMOUSKY, AND W.R. SORENSON, EFFECT OF COMPOSITION AND HIGH ENERGY RATE FORGING ON THE ONSET OF PRECIPITATION IN AN IRON-BASE SUPERALLOY, METALL. TRANS. A, VOL 13, MARCH 1982, P 345-353 16. "ARMCO 17-4PH STAINLESS STEEL BAR AND WIRE," PRODUCT DATA BULLETIN NO. S-24, ARMCO, INC., MIDDLETOWN, OH, SEPT 1966 17. "ARMCO PH 13-8MO STAINLESS STEEL," PRODUCT DATA BULLETIN NO. S-24, ARMCO, INC., MIDDLETOWN, OH, OCT 1986 18. "CARPENTER CUSTOM 450," PRODUCT DATA BULLETIN, CARPENTER TECHNOLOGY CORP., READING, PA, 1971 19. "CARPENTER CUSTOM 455," PRODUCT DATA BULLETIN, CARPENTER TECHNOLOGY CORP., READING, PA, 1971 20. "ARMCO 17-7PH AND PH15-7MO STAINLESS STEEL SHEET AND STRIP," PRODUCT DATA BULLETIN, ARMCO, INC., MIDDLETOWN, OH, JAN 1975 21. B. POLLARD, UNPUBLISHED WORK 22. N. SUUTALA, EFFECT OF SOLIDIFICATION CONDITIONS ON THE SOLIDIFICATION MODE IN

STAINLESS STEELS, METALL. TRANS. A, VOL 14, FEB 1983, P 191-197 23. G.E. LINNERT, WELDING PRECIPITATION-HARDENING STAINLESS STEELS, WELD. J., VOL 36, JAN 1957, P 9-27 24. F.J. HAWKINS, WELDING OF AGE-HARDENABLE STAINLESS STEELS, WELD. RES. COUNC. BULL., NO. 103, FEB 1965 25. "ARMCO FABRICATING DATA," BULLETIN SF-2, ARMCO, INC., MIDDLETOWN, OH, FEB 1966 26. C. CHEN AND J.S. YEH, ELECTRON BEAM AND PLASMA ARC WELDING OF 17-4PH STAINLESS STEEL, PROC. CONF. POWER BEAM PROCESSING: ELECTRON, LASER, PLASMA ARC (SAN DIEGO, CA), ASM INTERNATIONAL, 1988, P 45-54 27. A.J. TURNER, ELECTRON BEAM WELDING THICK SECTION PRECIPITATION-HARDENING STEEL, WELD. J., VOL 60, JAN 1981, P 18-26 28. M.U. ISLAM, G. CAMPBELL, AND R. HSU, FATIGUE AND TENSILE PROPERTIES OF EB WELDED 17-4PH STEEL, WELD. J., VOL 68, SEPT 1989, P 45-50 29. S. MATSUI ET AL., ELECTRON BEAM WELDING OF THE PRECIPITATION HARDENING STAINLESS STEEL, REVUE DE LA SOUDURE-LASTIJDSCRIFT, VOL 46, JAN 1990, P 52-57 30. W.R. CIESLAK, J.A. BROOKS, AND W.M. GARRISON, JR., THE WELDABILITY, MICROSTRUCTURE AND PROPERTIES OF PRECIPITATION STRENGTHENED MARTENSITIC STAINLESS STEELS, PROC. CONF. ADVANCES IN WELDING SCIENCE AND TECHNOLOGY (GATLINBURG, TN), ASM INTERNATIONAL, 1986 31. A.J. TURNER, ELECTRON BEAM WELDING OF HIGH VOLUME AIRCRAFT COMPONENTS, PROC. CONF. FUNDAMENTAL AND PRACTICAL APPROACHES TO THE RELIABILITY OF WELDED STRUCTURES (OSAKA, JAPAN), VOL 1, JAPAN WELDING SOCIETY, 1982, P 139-144 32. "WELDING CARPENTER CUSTOM 450 MULTIPLE-STRENGTH STAINLESS STEEL," PRODUCT DATA BULLETIN, CARPENTER TECHNOLOGY CORP., READING, PA, JAN 1973 33. ALLOY DIG., FILING CODE SS-193, APRIL 1967 34. R. KALTENHAUSER, WELDABILITY OF PRECIPITATION HARDENING STEELS, MET. ENG. QUART., VOL 9 (NO. 1), FEB 1969, P 44-57 35. "ARMCO FABRICATING DATA," BULLETIN SF-5, ARMCO, INC., MIDDLETOWN, OH, DEC 1967 36. "AM-350/AM-355 PRECIPITATION HARDENING STAINLESS STEELS," PRODUCT DATA BULLETIN, ALLEGHENY LUDLUM CORP., PITTSBURGH, PA, 1963 37. H.W. MISHLER, R.E. MONROE, AND P.J. RIEPPEL, "WELDING OF HIGH STRENGTH STEELS FOR AIRCRAFT AND MISSILE APPLICATIONS," DMIC REPORT 118, DEFENSE METALS INFORMATION CENTER, COLUMBUS, OH, 12 OCT 1959 38. J.J. VAGI AND D.C. MARTIN, WELDING OF HIGH-STRENGTH STAINLESS STEELS FOR ELEVATED TEMPERATURE USE, WELD. J., VOL 35, MARCH 1956, P 137S-144S 39. B.S. BLUM AND R.H. WITT, HEAT AFFECTED ZONE CRACKING IN A-286 WELDMENTS, WELD. J., VOL 42, AUG 1963, P 365S-370S 40. J.A. BROOKS AND R.W. KRENZER, PROGRESS TOWARD A MORE WELDABLE A-286, WELD. J., VOL 53, JUNE 1974, P 242S-245S 41. J.A. BROOKS, EFFECT OF ALLOY MODIFICATIONS ON HAZ CRACKING OF A-286 STAINLESS STEEL, WELD. J., VOL 53, NOV 1974, P 517S-523S 42. J.A. BROOKS, A.W. THOMPSON, AND J.C. WILLIAMS, WELD CRACKING OF AUSTENITIC STAINLESS STEELS--EFFECTS OF IMPURITIES AND MINOR ELEMENTS, PROC. CONF. PHYS. METALLURGY OF METAL JOINING (ST. LOUIS, MO), TMS/AIME, 1980, P 117-136 43. W.H. SIMON, S. BANERJEE, AND R.S. CREMISIO, METALLURGICALLY CONTROLLED WELDING FILLER WIRE--TRACE ELEMENT CONTROLLED TO PRODUCE CONSISTENTLY SOUND WELDS, PROC. INT. CONF.: THE EFFECTS OF RESIDUAL IMPURITY AND MICROALLOYING ELEMENTS ON WELDABILITY AND WELD PROPERTIES (LONDON, ENGLAND), WELDING INSTITUTE, PAPER 14, 1983

44. D.W. HARVEY, METALLURGICALLY CONTROLLED WELDING FILLER WIRE, WELD. REV., MAY 1988, P 80, 82, 122, 123 45. L.T. SUMMERS, "EFFECT OF CHEMICAL SEGREGATION ON THE MICROSTRUCTURE AND PROPERTIES OF A-286 WELDMENTS," PH.D. DISSERTATION, UNIVERSITY OF CALIFORNIA, BERKELEY, 1984 46. A. INOUE ET AL., FATIGUE STRENGTH OF ALLOY A-286 AT CRYOGENIC TEMPERATURE, PROC. CONF. FATIGUE '87, VOL 1 (CHARLOTTESVILLE, VA), ENGINEERING MATERIALS ADVISORY SERVICES, LTD., 1987, P 547-555 47. R.E. ROTH AND N.F. BRATKOVICH, CHARACTERISTICS AND STRENGTH DATA OF ELECTRON BEAM WELDS IN FOUR REPRESENTATIVE MATERIALS, WELD. J., VOL 41, MAY 1962, P 229S-240S 48. I.F. SQUIRES, DETERMINATION OF RESISTANCE SPOT WELDING CONDITIONS FOR SOME HEAT-RESISTING STAINLESS STEELS, BR. WELD. J., VOL 9, MARCH 1962, P 149-157 49. P.M. KNOWLSON, DETERMINATION OF RESISTANCE SEAM WELDING CONDITIONS FOR SOME HEAT RESISTING STEELS, BR. WELD. J., VOL 9, 1962, P 436-445 50. J.J. VAGI, R.M. EVANS, AND D.C. MARTIN, "WELDING PRECIPITATION-HARDENING STAINLESS STEELS," NASA TECHNICAL MEMORANDUM X-53582, 28 FEB 1967 51. B.I. MEDOVAR, ON THE NATURE OF WELD HOT CRACKING, TRANSLATED FROM AVTOM. SVARKA, VOL 7 (NO. 4), 1954, P 12-28 52. M.J. STRUM, L.T. SUMMERS, AND J.W. MORRIS, JR., THE AGING RESPONSE OF A WELDED IRON-BASED SUPERALLOY, WELD. J., VOL 62, SEPT 1983, P 235S-242S Selection of Cast Stainless Steels Michael J. Cieslak, Sandia National Laboratories

Introduction THE DEVELOPMENT OF CASTING TECHNOLOGY has been driven by two distinct engineering perspectives. The first involves the potential economy of cast hardware relative to totally machined wrought parts. Complex, nearnet shapes can be produced by a variety of casting processes, including sand casting, investment casting, and centrifugal casting (used almost exclusively for making tubular products). Casting also can be an extremely effective way of producing very large parts. The second perspective is associated with metallurgical realities. Many alloys have been developed for which there are no true wrought counterparts. The chemical composition and microstructure of these alloys is such that deformation processing of large ingots into wrought product forms would be problematic. In addition, the very nature of the cast microstructure, which in general is composed of large grains, can provide benefits in high-temperature creep applications. Some common examples of stainless steel castings include valve and pump bodies, compressor wheels, propellers, turbine impellers, large pipe elbows, heat treating fixtures, and food, paper, and chemical processing equipment. Extensive reviews on the properties and applications of cast stainless steels can be found in Ref 1, 2, 3, and 4. Acknowledgement This work was performed at Sandia National Laboratories, supported by the U.S. Department of Energy under contract number DE-AC04-76DP00789.

References

1. J.M. SVOBODA, HIGH-ALLOY STEELS, CASTING, 9TH ED., VOL 15, METALS HANDBOOK, ASM INTERNATIONAL, 1988, P 722-735 2. M. BLAIR, CAST STAINLESS STEELS, PROPERTIES AND SELECTION: IRONS, STEELS, AND HIGHPERFORMANCE ALLOYS, VOL 1, ASM HANDBOOK, ASM INTERNATIONAL, 1990, P 908-929 3. R.W. MONROE AND S.J. PAWEL, CORROSION OF CAST STEELS, CORROSION, 9TH ED., VOL 13, METALS HANDBOOK, ASM INTERNATIONAL, 1987, P 573-582 4. P.R. WEISER, ED., STEEL CASTINGS HANDBOOK, 5TH ED., STEEL FOUNDERS' SOCIETY OF AMERICA, 1980 Selection of Cast Stainless Steels Michael J. Cieslak, Sandia National Laboratories

Stainless Steel Casting Alloys Categories. Unlike the case of wrought alloy designators, cast stainless steels are not grouped according to microstructural types, but by the intended service environment. The general classification of stainless steel castings places these alloys in one of two categories: corrosion-resistant alloys and heat-resistant alloys. The corrosion-resistant alloys have been designed primarily for service environments involving aqueous or liquid-vapor corrosives at temperatures generally less than 315 °C (600 °F). Use at extended temperatures up to 650 °C (1200 °F) in selected environments may also be considered.

The heat-resistant alloys have been designed for use in service environments exceeding 650 °C (1200 °F) where aqueous corrosives do not exist. At these elevated temperatures, oxidation, sulfidation, and carburization/decarburization are the environmental effects of consequence. High-temperature mechanical strength, creep and stress-rupture resistance, and microstructural stability are the engineering requirements for these alloys. The high-temperature alloys are most readily distinguished from the corrosion-resistant alloys on the basis of carbon content (although other alloying elements may also vary considerably in response to service needs) with the high-temperature alloys containing 0.20 to 0.70 wt% C and the corrosion-resistant alloys generally containing less than 0.20 wt% C. Nomenclature. Cast stainless steels are most often specified on the basis of composition using the designation system of

the High Alloy Product Group of the Steel Founders' Society of America. (The High Alloy Product Group has replaced the Alloy Casting Institute, or ACI, which formerly administered these designations.) The first letter of the designation indicates whether the alloy is intended primarily for liquid corrosion service (C) or high-temperature service (H). The second letter in either group of alloys indicates an approximate ratio of nickel-to-chromium with the ratio increasing in alphabetical order. (As the nickel content increases, the second letter of the designation progresses through the alphabet from A to Z.) The number immediately following the hyphen indicates the carbon content of the alloy in hundredths of a weight percent (also known as "points" of carbon). For the corrosion-resistant alloys, the carbon content indicated is the maximum allowable for the alloy. For the heat-resistant alloys, the carbon content identified is the midpoint with an allowable range of ± 0.05 wt%. Finally, if further alloying elements are present, these are indicated by the addition of one or more letters as a suffix. Thus, the designation of CF-8M refers to an alloy for corrosion-resistant service (C) of the 19Cr-9Ni type with a maximum carbon content of 0.08% that contains molybdenum (M). Figure 1 graphically shows the range of nickel and chromium contents in several grades of corrosion-resistant and heat-resistant alloys identified by the aforementioned designation system. Tables 1 and 2 list the chemical compositions of the common corrosion-resistant and heat-resistant alloys, respectively.

TABLE 1 COMPOSITIONS AND TYPICAL MICROSTRUCTURES OF CORROSION-RESISTANT STAINLESS STEEL CASTING ALLOYS

ALLOY WROUGHT DESIGNATION ALLOY TYPE(A) CHROMIUM STEELS CA-15 410

ASTM SPECIFICATIONS

MOST COMMON END-USE MICROSTRUCTURE

COMPOSITIONS, %(B) C Mn Si Cr

Ni

OTHERS(C)

A 743, A 217, A 487

MARTENSITE

0.15

1.00

1.50

1.0

0.50 MO(D)

CA-15M

...

A 743

MARTENSITE

0.15

1.00

0.65

1.0

0.15-1.00 MO

CA-40

420

A 743

MARTENSITE

0.40

1.00

1.50

1.0

0.5 MO(D)

CA-40F

...

A 743

MARTENSITE

1.00

1.50

1.0

...

CB-30

431, 442

A 743

1.00

1.50

2.0

...

CC-50

446

A 743

FERRITE AND CARBIDES FERRITE AND CARBIDES

0.20.4 0.30 0.30

1.00

1.50

11.514.0 11.514.0 11.514.0 11.514.0 18.022.0 26.030.0

4.0

...

1.012.5 11.514.0 11.012.5 15.517.7

6.08.0 3.54.5 0.501.00 3.64.6

...

CHROMIUM-NICKEL STEELS CA-6N ... A 743

MARTENSITE

0.06

0.50

1.00

CA-6NM

...

A 743, A 487

MARTENSITE

0.06

1.00

1.00

CA-28MWV

...

A 743

MARTENSITE

...

A 747

MARTENSITE, AGE HARDENABLE

0.501.00 0.70

1.00

CB-7CU-1

0.200.28 0.07

CB-7CU-2

...

A 747

MARTENSITE, AGE HARDENABLE

0.07

0.70

1.00

14.015.5

4.55.5

CD-4MCU

...

1.00

1.00

312

AUSTENITE IN FERRITE, AGE HARDENABLE FERRITE IN AUSTENITE

0.04

CE-30

A 351, A 743, A 744, A 890 A 743

0.30

1.50

2.00

CF-3(E)

304L

A 351, A 743, A 744

FERRITE IN AUSTENITE

0.03

1.50

2.00

25.026.5 26.030.0 17.0-

4.756.0 8.011.0 8.0-

1.00

0.4-1.0 MO 0.9-1.25 MO; 0.91.25 W; 0.2-0.3 V 2.5-3.2 CU; 0.200.35 NB; 0.05 N MAX 2.5-3.2 CU; 0.200.35 NB; 0.05 N MAX 1.75-2.25 MO; 2.75-3.25 CU ... ...

(E)

CF-3M

316L

A 351, A 743, A 744


0.03

1.50

2.00

CF-3MN

...

A 743

FERRITE TN AUSTENITE

0.03

1.50

1.50

CF-8(E)

304

A 351, A 743, A 744


0.08

1.50

2.00

CF-8C

347

A 351, A 743, A 744


0.08

1.50

2.00

CF-8M

316

A 351, A 743, A 744


0.08

1.50

2.00

CF-10

...

A 351


1.50

2.00

CF-10M

...

A 351


1.50

1.50

CF-10MC

...

A 351


0.040.10 0.040.10 0.10

1.50

1.50

CF-10SMNN

...

A 351, A 743


0.10

CF-12M

316

...

0.12

3.504.50 2.00

CF-16F

303

A 743

FERRITE IN AUSTENITE OR AUSTENITE AUSTENITE

7.009.00 1.50

0.16

1.50

2.00

CF-20

302

A 743

AUSTENITE

0.20

1.50

2.00

CG-6MMN

...

A 351, A 743


0.06

4.006.00

1.00

CG-8M

317

A 351, A 743, A 744


0.08

1.50

1.50

CG-12

...

A 743


0.12

1.50

2.00

CH-8

...

A 351


0.08

1.50

1.50

CH-10

...

A 351


0.040.10

1.50

2.00

21.0 17.021.0 17.021.0 18.021.0 18.021.0 18.021.0 18.021.0 18.021.0 15.018.0 16.018.0 18.021.0 18.021.0 18.021.0 20.523.5

12.0 8.012.0 9.013.0 8.011.0 9.012.0 9.012.0 8.011.0 9.012.0 13.016.0 8.09.0 9.012.0 9.012.0 8.011.0 11.513.5

18.021.0 20.023.0 22.026.0 22.026.0

9.013.0 10.013.0 12.015.0 12.015.0

2.0-3.0 MO 2.0-3.0 MO; 0.100.20 N ... NB(F) 2.0-3.0 MO ... 2.0-3.0 MO 1.75-2.25 MO 0.08-0.18 N 2.0-3.0 MO 1.50 MO MAX; 0.20-0.35 SE ... 1.50-3.00 MO; 0.10-0.30 NB; 0.10-30 V; 0.20-40 N 3.0-4.0 MO ... ... ...

CH-20

309

A 351, A 743

AUSTENITE

0.20

1.50

2.00

22.026.0 19.520.5

12.015.0 17.519.5

CK-3MCUN

...

A 351, A 743, A 744


0.025

1.20

1.00

CK-20

310

A 743

AUSTENITE

0.20

2.00

2.00

23.027.0

19.022.0

NICKEL-CHROMIUM STEEL CN-3M ... A 743

AUSTENITE

0.03

2.00

1.00

0.07

1.50

1.50

AUSTENITE

0.07

1.50

3.50(G)

AUSTENITE

0.050.15

0.151.50

0.501.50

20.022.0 19.022.0 18.020.0 19.021.0

23.027.0 27.530.5 22.025.0 31.034.0

CN-7M

...

A 351, A 743, A 744

AUSTENITE

CN-7MS

...

A 743, A 744

CT-15C

...

A 351

... 6.0-7.0 V; 0.180.24 N; 0.50-1.00 CU ...

4.5-5.5 MO 2.0-3.0 MO; 3.04.0 CU 2.5-3.0 MO; 1.52.0 CU 0.5-1.5 V

(A) TYPE NUMBERS OF WROUGHT ALLOYS ARE LISTED ONLY FOR NOMINAL IDENTIFICATION OF CORRESPONDING WROUGHT AND CAST GRADES. COMPOSITION RANGES OF CAST ALLOYS ARE NOT THE SAME AS FOR CORRESPONDING WROUGHT ALLOYS; CAST ALLOY DESIGNATIONS SHOULD BE USED FOR CASTINGS ONLY. (B) MAXIMUM UNLESS A RANGE IS GIVEN. THE BALANCE OF ALL COMPOSITIONS IS IRON. (C) SULFUR CONTENT IS 0.04% IN ALL GRADES EXCEPT: CG-6MMN, 0.030% S (MAX); CF-10SMNN, 0.03% S (MAX); CT-15C, 0.03% S (MAX); CK-3MCUN, 0.010% S (MAX); CN-3M 0.030% S (MAX), CA-6N, 0.020% S (MAX); CA-28MWV, 0.030% S (MAX); CA-40F, 0.20-0.40% S; CB-7CU-1 AND -2, 0.03% S (MAX). PHOSPHORUS CONTENT IS 0.04% (MAX) IN ALL GRADES EXCEPT; CF-16F, 0.17% P (MAX); CF10SMNN, 0.060% P (MAX); CT-15C, 0.030% P (MAX); CK-3MCUN, 0.045% P (MAX); CN-3M, 0.030% P (MAX); CA-6N, 0.020% P (MAX); CA28MWV, 0.030% P (MAX); CB-7CU-1 AND -2, 0.035% P (MAX). (D) MOLYBDENUM NOT INTENTIONALLY ADDED. (E) CF-3A, CF-3MA, AND CF-8A HAVE THE SAME COMPOSITION RANGES AS CF-3, CF-3M, AND CF-8, RESPECTIVELY, BUT HAVE BALANCED COMPOSITIONS SO THAT FERRITE CONTENTS ARE AT LEVELS THAT PERMIT HIGHER MECHANICAL PROPERTY SPECIFICATIONS THAN THOSE FOR RELATED GRADES. THEY ARE COVERED BY ASTM A 351. (F) NB, 8 × %C MIN (1.0% MAX ); OR NB + TA × %C (1.1% MAX). (G) FOR CN-7MS, SILICON RANGES FROM 2.50 TO 3.50%. TABLE 2 COMPOSITIONS OF HEAT-RESISTANT STAINLESS STEEL CASTING ALLOYS

ALLOY

UNS

ASTM SPECIFICATIONS(A) COMPOSITION, %(B)

DESIGNATION HA HC HD HE HF HH HI HK HK30 HK40 HL HN HP HP-50WZ(C) HT HT30 HU HW HX

NO. ... J92605 J93005 J93403 J92603 J93503 J94003 J94224 ... ... J94604 J94213 ... ... J94605 ... ... ... ...

A217 A 297, A 608 A 297, A 608 A 297, A 608 A 297, A 608 A 297, A 608, A 447 A 297, A 567, A 608 A 297, A 351, A 567, A 608 A 351 A 351 A 297, A 608 A 297, A 608 A 297 ... A 297, A 351, A 567, A 608 A 351 A 297, A 608 A 297, A 608 A 297, A 608

C 0.20 MAX 0.50 MAX 0.50 MAX 0.20-0.50 0.20-0.40 0.20-0.50 0.20-0.50 0.20-0.60 0.25-0.35 0.35-0.45 0.20-0.60 0.20-0.50 0.35-0.75 0.45-0.55 0.35-0.75 0.25-0.35 0.35-0.75 0.35-0.75 0.35-0.75

Cr 8-10 26-30 26-30 26-30 19-23 24-28 26-30 24-28 23.0-27.0 23.0-27.0 28-32 19-23 24-28 24-28 13-17 13.0-17.0 17-21 10-14 15-19

Ni ... 4 MAX 4-7 8-11 9-12 11-14 14-18 18-22 19.0-22.0 19.0-22.0 18-22 23-27 33-37 33-37 33-37 33.0-37.0 37-41 58-62 64-68

Si (max) 1.00 2.00 2.00 2.00 2.00 2.00 2.00 2.00 1.75 1.75 2.00 2.00 2.00 2.50 2.50 2.50 2.50 2.50 2.50

(A) ASTM DESIGNATIONS ARE THE SAME AS ACI DESIGNATIONS. (B) BALANCE IRON IN ALL COMPOSITIONS. MANGANESE CONTENT: 0.35 TO 0.65% FOR HA. 1% FOR HC. 1.5% FOR HD AND 2% FOR THE OTHER ALLOYS. PHOSPHORUS AND SULFUR CONTENTS: 0.04% (MAX) FOR ALL BUT HP-50WZ. MOLYBDENUM IS INTENTIONALLY ADDED ONLY TO HA, WHICH HAS 0.90 TO 1.20% MO; MAXIMUM FOR OTHER ALLOYS IS SET AT 0.5% MO. HH ALSO CONTAINS 0.2% N (MAX). (C) ALSO CONTAINS 4 TO 6% W, 0.1 TO 1.0% ZR, AND 0.035% S (MAX) AND P (MAX)

FIG. 1 SCHEMATIC OF THE CHROMIUM AND NICKEL CONTENT RANGES FOR CORROSION-RESISTANT AND HEAT-RESISTANT STAINLESS STEEL ALLOYS. THE LETTERS IN CIRCLES REPRESENT THE SECOND LETTER IN THE ALLOY DESIGNATION SYSTEM. SOURCE: REF 4


4. P.R. WEISER, ED., STEEL CASTINGS HANDBOOK, 5TH ED., STEEL FOUNDERS' SOCIETY OF AMERICA, 1980 Selection of Cast Stainless Steels

Michael J. Cieslak, Sandia National Laboratories

Welding and Weldability of Cast Stainless Steels Welding and Welding Processes. In general terms, there are two situations in which welding of stainless steel castings

is required. One of these occurs before the casting ever leaves the foundry. Casting defects, such as hot tears, shrinkage cavities, and cold shuts, are often observed in stainless steel castings. Certain specifications preclude the acceptance of these castings for service use without repair. Fusion welding is an appropriate method for repair of these casting defects. The defects, identified through one of several possible procedures (for example, dye penetrant, radiography, or ultrasonic inspection), are removed by machining or grinding, leaving a clean joint surface onto which weld filler metal is added to produce a sound final product. For small, near-surface defects, autogenous welds (that is, no filler metal added) may be sufficient to eliminate the defects. A special case of repair welding is known by the term "upgrading." An "upgraded" casting is one that has been selectively welded in order to achieve a higher level of product specification acceptance. The second situation requiring welding involves joining of the casting to other hardware, cast or wrought, as part of an engineering design. In this circumstance, the design rules for welding are similar to those for wrought alloys, as discussed in the articles on wrought stainless steels that appear in this Section of the Handbook. Any of the fusion welding processes can be successfully employed in the welding of castings. Most common are the arc welding processes, especially autogenous and filler-metal-added gas-tungsten arc welding, shielded metal arc welding, and gas-metal arc welding. Fabrication involving the use of extremely large castings may employ electroslag welding as the only reasonable economic choice. Other processes, such as electron-beam welding, laser-beam welding, and plasma arc welding, are also appropriate choices in many instances. A wide range of filler metals exists for joining of the various casting alloys. The choice depends on the specifics of the situation and whether the weld will be to a similar or dissimilar alloy. The use of matching filler metal is generally suggested when welding like alloys. Where a matching consumable is not available, common choices are 308L and 316L for the low-nickel-content corrosion-resistant austenitic alloys. For the highernickel-content corrosion-resistant and heat-resistant alloys, appropriate choices may be consumables (bare electrodes for gas-tungsten arc welding and gas-metal arc welding) such as 20Cb-3 (AWS ER320/ER320LR) and nickel alloys such as C-276 (AWS ERNiCrMo-4), Alloy 625 (ERNiCrMo-3), and Inconel alloys 82 (AWS ERNiCr-3), 92 (AWS ERNiCrFe6), and 182 (AWS ERNiCrFe-3). Metallurgical Considerations. Castings are metallurgically more complex than wrought alloys. Segregation during

dendritic solidification can lead to local variations in both chemical composition and microstructure. The vast majority of the heat-resistant alloys will terminate solidification with the formation of a eutectic-like carbide constituent that remains upon cooling of the casting to room temperature. This microstructural constituent helps to augment the creep resistance of this group of alloys. Subsequent heat treatment of these alloys is not common. Many of the lower-nickel-content corrosion-resistant alloys (for example, CF-3, CF-8M, and CB-7Cu) will have a duplex microstructure consisting of ferrite islands in a matrix of either austenite or martensite. Postcasting solution heat treatments used to prepare corrosion-resistant alloys for service are insufficient to homogenize the microstructure. The inhomogeneous microstructure of both groups of alloys will tend to promote relatively wide areas of partial melting adjacent to the fusion zone when compared with similar wrought alloys. Martensitic Stainless Steel Castings. Special considerations exist for the welding of martensitic stainless steel

castings (CA-6NM, CA-15, CA40, and CB-7Cu-1, and CB-7Cu-2). Quench cracking and a marked reduction in mechanical properties (especially ductility and impact toughness) may result from the welding of these alloys. Corrosion resistance may also be detrimentally affected by the thermal cycle that the material undergoes in the fusion zone and heataffected-zone (HAZ) of welds. Suggested welding practice often calls for the preheating and postheating of most of the martensitic alloys. Alloys CA-15 and CA-40 require preheating temperatures in the range of 200 to 315 °C (390 to 600 °F) and a postweld heat treatment (PWHT) temperature range of 610 to 760 °C (1130 to 1400 °F). Alloy CA-6NM requires a preheat temperature range of 100 to 150 °C (210 to 300 °F) and a PWHT temperature range of 590 to 620 °C (1095 to 1150 °F). The higher preheating temperatures should be used when welding thicker sections and higher carbon compositions of these alloys.

The PWHT for these alloys is a tempering operation intended to restore ductility and toughness at the expense of strength and hardness. Because of the broad range of potential tempering temperatures, a wide range of PWHT properties are possible. Alloy CB-7Cu (1 and 2) has no suggested preheat temperature, but a range (480 to 590 °C, or 895 to 1095 °F) of PWHT temperatures is recommended. In the case of this alloy, the PWHT is an aging operation. In order to recover properties in the welded product, a full solution anneal (1040 °C, or 1095 °F, oil quench), followed by aging, is required. Direct aging without solution annealing will not fully restore preweld properties. The higher aging temperatures result in the lowest strength but also provide the highest ductility and toughness. Welding Defects. The two most serious problems encountered in the welding of stainless steel castings, either during the repair of casting defects or for subsequent attachment to other structures, are solidification hot cracking and heataffected zone (HAZ) hot cracking. Figure 2 shows a prototypical hot crack extending both into the fusion zone and the HAZ of a casting alloy. Both of these phenomena involve separation of grain boundaries due to the presence of a wetting liquid phase and sufficient mechanical imposition.

FIG. 2 OPTICAL MICROGRAPH OF THE FUSION LINE REGION IN A CAST CORROSION-RESISTANT AUSTENITIC STAINLESS STEEL. NOTE THE HOT CRACK EXTENDING BOTH INTO THE FUSION ZONE (LEFT) AND BACK INTO THE PARTIALLY MELTED REGION OF THE HAZ. ARROWS INDICATE REGIONS OF PARTIAL MELTING.

Studies (Ref 5, 6, 7, 8, 9) have indicated that wholly austenitic stainless steel casting alloys are more susceptible to the formation of these types of defects than are duplex ferrite plus austenite alloys. In particular, Cieslak (Ref 7, 8) and Cadden (Ref 9) have shown that CF-8M alloys (being prototypic of a wide range of low-nickel corrosion-resistant casting alloys) having welds that solidify with ferrite as the primary solidification phase are much more resistant to hot cracking than are alloys that solidify with austenite as the primary solidification phase. Table 3 lists the chemical compositions and ferrite content of a series of CF-8M castings used in a study of weldability (Ref 8). Figure 3 shows the fusion zone hot cracking susceptibility of these alloys determined using the Varestraint test at 1 % applied strain. The alloys that solidified as primary austenite (γ) were more susceptible to cracking than those which 2

solidified as primary ferrite (δ). With the exception of heat number 7, the primary ferrite alloys did not experience any cracking under this set of experimental conditions. The high level of sulfur (0.032 wt%) and phosphorus (0.046 wt%) were identified as being responsible for the greater cracking sensitivity in heat number 7. This particular observation was important for establishing that ferrite content alone was not a sufficient criterion to preclude the possibility of hot cracking in duplex alloys.

TABLE 3 COMPOSITIONS OF ALLOYS IN CF-8M WELDABILITY STUDY

HEAT NO. 1 2 3 4

COMPOSITION, WT%(A) C Mn Si Cr Ni 0.08 0.60 1.05 18.32 13.20 0.06 0.44 1.08 21.10 9.55 0.05 0.26 0.83 17.65 12.03 0.06 1.17 0.60 18.18 12.08

FERRITE, VOL% Mo S P N 2.26 0.016 0.035 0.041 0.17 2.52 0.008 0.021 0.048 22.8 2.05 0.017 0.021 0.06 1.48 2.48 0.016 0.027 0.05 4.16

5 6 7

0.07 0.57 0.99 18.21 9.57 2.39 0.019 0.024 0.06 0.04 0.21 0.38 19.55 15.38 2.88 0.025 0.036 0.04 0.10 0.30 0.69 21.31 10.62 2.34 0.032 0.046 0.05

7.60 2.00 14.24

Source: Ref 8

(A)

THE BALANCE IS IRON.

FIG. 3 VARESTRAINT TEST RESULTS ON SEVERAL HEATS OF CF-8M CORROSION-RESISTANT ALLOY. TESTS WERE RUN AT

1 % AUGMENTED STRAIN USING AN ENERGY INPUT OF APPROXIMATELY 2.7 KJ/MM (69 2

KJ/IN.). SEE TABLE 3 FOR HEAT NUMBER CHEMICAL COMPOSITIONS.

In the same investigation, the hot cracking sensitivity of the heat-resistant alloy HK-40 was also examined. This alloy solidifies as primary austenite and terminates solidification with a eutectic-like Cr7C3 carbide constituent. This is prototypical solidification behavior of a wide range of heat-resistant stainless steel casting alloys. It was found that cracking was exacerbated by low carbon content (0.38 wt%), and high levels of combined sulfur and phosphorus (>0.026 wt%). Silicon addition also had a detrimental effect on cracking susceptibility when either carbon content was low or sulfur or phosphorus was high. It was observed that hot cracks could be "healed" by the carbide eutecticlike constituent, suggesting that higher carbon contents in alloys of this type could be helpful in minimizing this type of defect. Similar results have been found for 16Cr-35Ni (HT type) (Ref 10) and high-carbon-content alloys such as 25Cr-20Ni (the cast equivalent of wrought type 310 containing up to 0.2 wt% C) (Ref 11). A study of the weldability of alloy CN-7M identified sulfur, phosphorus, and silicon as being detrimental to the hot cracking resistance of this high-nickel-content (29 wt% Ni) corrosion-resistant alloy (Ref 12). Specifically, silicon was found to promote the formation of an M6C carbide eutecticlike constituent associated with intergranular hot cracked regions. For a discussion of the hot cracking susceptibility of a wide range of corrosion-resistant and heat-resistant casting alloys, along with microstructural analysis of the cracking behavior, a review of the work of Cadden (Ref 9) is suggested. In many studies examining both the fusion zone and HAZ hot cracking behavior of stainless steel casting alloys, the presence of sulfur and phosphorus has universally been identified with increased cracking susceptibility. To a somewhat lesser extent, so have silicon and niobium. Carbon can also be detrimental, although not all studies have reached the same conclusion. It can be said that if there is sufficient carbon present to result in the formation of terminal-stage, eutectic-like constituents, then the hot cracking susceptibility will increase up to the point where the level of carbon promotes eutectic healing. Eutectic healing may eliminate the hot cracking problem, but the reduction in ductility with increasing carbon content can limit the utility of the weld joint for an engineering application (Ref 10). Each of these elements appears to increase cracking susceptibility in two possible ways. The first is by promoting the formation of low-melting-point constituents, and the second is by acting as a surfactant which promotes wetting of solidifying grain boundaries. As discussed elsewhere in this Volume (see, for example, the articles in this Section of the

Handbook that describe wrought stainless steels), either phenomenon will promote an increased tendency toward fusion zone or HAZ hot cracking. A specific concern in the welding of corrosion-resistant stainless steel castings is the production of a sensitized microstructure. Sensitization occurs in stainless steel alloys when chromium-rich phases--M23C6 carbides, sigma (σ), chi (χ), and Laves phases--precipitate during exposure to temperatures between 400 and 900 °C (750 and 1650 °F). The precipitation of these constituents can result in a local depletion of chromium (and also molybdenum), lowering the corrosion resistance to a point where corrosion-induced failure of the casting can result. Primary among the phases of concern is the M23C6 carbide (with M being primarily chromium). Precipitation of carbides at grain boundaries can occur during the fusion welding cycle at sites in both the HAZ and the fusion zone. Because the kinetics of precipitation are proportional to the carbon concentration, low-carbon grades (such as CF-3 or CF-3M) have been developed that can experience welding without producing a sensitized microstructure. A final metallurgical consideration involves the welding of castings that have seen extended service at relatively high temperatures. For a variety of reasons, castings may have to be rewelded sometime during their service lifetime. This can pose a difficult problem if the temperature environment has modified the alloy microstructure so that additional constituents are present that adversely affect alloy ductility. It is always preferable to heat treat the casting to remove these unwanted microstructural features. This process may require solutionizing at temperatures above 1000 °C (1830 °F) and may be impractical for large castings. If solution heat treatment is not an option, then special care must be taken to ensure that the new weld joint design will have the lowest possible level of inherent restraint. Weld Parameters. From an understanding of the metallurgical behavior of stainless steel casting alloys given above,

general rules for fusion welding can be inferred. Low-heat-input processes are almost universally suggested. For alloys that are far from homogeneous, a high-heat-input process would produce flatter thermal gradients outside of the fusion zone, creating a wider zone of partial melting and a greater microstructural sensitivity to HAZ hot cracking. The greater total shrinkage strain associated with the solidification of a large (high-energy-input) weld can create a restraint situation more likely to result in fusion zone solidification cracking. The flatter thermal gradients of higher-energy-input welds reduce the cooling rates through the carbide precipitation temperature range, resulting in a higher likelihood of a sensitized structure in the corrosion-resistant alloys. For the same reasons, low interpass temperatures are generally employed in the welding of these alloys. Interpass temperatures below 200 °C (400 °F) are common. For critical applications (especially among corrosion-resistant alloys), interpass temperatures of less than 150 °C (300 °F) are often specified. Spraying castings with water between passes is not uncommon for the austenitic corrosion-resistant alloys (Ref 4). Preheating of castings prior to welding is almost never required and generally not suggested, at least for the corrosionresistant alloys, because the sensitization impact of such processing is similar to the use of high-energy-input welding parameters. In the case of heat-resistant alloy castings, preheating can be used to reduce the magnitude of thermal shrinkage strains, but again this benefit comes at the expense of a wider zone of partial melting. There is no known formula for preheating temperature and welding parameters for a given alloy that can ensure the preclusion of welding defects. Because of this technological limitation, the use of preheating prior to welding of these alloys is questionable. Postweld solutionizing of corrosion-resistant alloys is always recommended. For the heat treatable alloys (for example, CB-7Cu), a complete solutionizing and aging heat treatment is recommended to reestablish baseline properties. Where PWHT is not possible, it is even more important to weld using low-heat-input parameters. Where direct fusion welding is not appropriate, as in the case of an extremely hot-crack-sensitive casting, other approaches might prove useful. One option is to "butter" the area of the crack-sensitive casting to be joined, using a consumable of a less crack-sensitive material. This "buttering" involves surfacing the casting locally with a consumable before fabricating the actual weld joint. With no joint present during the buttering operation, the restraint on the solidifying surface pass and the underlying HAZ of the casting can be quite small and defects (fusion zone and HAZ hot cracks) may be avoided. This surfaced area can be built up slowly to a thickness where it can then be prepared as a joint for subsequent welded attachment. Additional information on buttering can be found in the article "Hardfacing, Weld Cladding, and Dissimilar Metal Joining" in this Volume. As another option, an insert of a less hot-crack-sensitive alloy can be inertia or friction welded onto the mating surface of the hot-crack-sensitive casting. This insert could then be prepared for subsequent fusion welding to the required structure. These options are contingent upon the availability of consumable materials that are acceptable in the service environment.


4. P.R. WEISER, ED., STEEL CASTINGS HANDBOOK, 5TH ED., STEEL FOUNDERS' SOCIETY OF AMERICA, 1980 5. H.S. AVERY, CAST HEAT-RESISTANT ALLOYS FOR HIGH-TEMPERATURE SERVICE, WELD. RES. COUNC. BULL., NO. 143, AUG 1969 6. "THE WELDING OF CAST AUSTENITIC STEELS," WELDING INSTITUTE REPORT SERIES, THE WELDING INSTITUTE, SEPT 1970, PHASES 1-4 7. M.J. CIESLAK AND W.F. SAVAGE, WELD. J., MAY 1980, P 136-S TO 146-S 8. M.J. CIESLAK AND W.F. SAVAGE, "WELDABILITY AND SOLIDIFICATION PHENOMENA OF CAST HIGH ALLOY HEAT RESISTANT ALLOY HK-40 AND CORROSION RESISTANT 19CR-9NI2.5MO STEELS," REPORT A-69, STEEL FOUNDERS' SOCIETY OF AMERICA, 1980 9. C.H. CADDEN, "WELDABILITY OF HIGH ALLOY CASTINGS," M.S. THESIS, RENSSELAER POLYTECHNIC INSTITUTE, TROY, NEW YORK, MAY 1981 10. D. ROZET, H.C. CAMPBELL, AND R.D. THOMAS, WELD. J., OCT 1949, P 481-S TO 491-S 11. Y. ARATA, F. MATSUDA, AND S. KATAYAMA, TRANS. JWRI, JUNE 1977, P 105-116 12. M.J. CIESLAK AND W.F. SAVAGE, WELD. J., MAY 1985, P 119-S TO 126-S Selection of Cast Stainless Steels Michael J. Cieslak, Sandia National Laboratories

References

1. J.M. SVOBODA, HIGH-ALLOY STEELS, CASTING, 9TH ED., VOL 15, METALS HANDBOOK, ASM INTERNATIONAL, 1988, P 722-735 2. M. BLAIR, CAST STAINLESS STEELS, PROPERTIES AND SELECTION: IRONS, STEELS, AND HIGH-PERFORMANCE ALLOYS, VOL 1, ASM HANDBOOK, ASM INTERNATIONAL, 1990, P 908-929 3. R.W. MONROE AND S.J. PAWEL, CORROSION OF CAST STEELS, CORROSION, 9TH ED., VOL 13, METALS HANDBOOK, ASM INTERNATIONAL, 1987, P 573-582 4. P.R. WEISER, ED., STEEL CASTINGS HANDBOOK, 5TH ED., STEEL FOUNDERS' SOCIETY OF AMERICA, 1980 5. H.S. AVERY, CAST HEAT-RESISTANT ALLOYS FOR HIGH-TEMPERATURE SERVICE, WELD. RES. COUNC. BULL., NO. 143, AUG 1969 6. "THE WELDING OF CAST AUSTENITIC STEELS," WELDING INSTITUTE REPORT SERIES, THE WELDING INSTITUTE, SEPT 1970, PHASES 1-4 7. M.J. CIESLAK AND W.F. SAVAGE, WELD. J., MAY 1980, P 136-S TO 146-S 8. M.J. CIESLAK AND W.F. SAVAGE, "WELDABILITY AND SOLIDIFICATION PHENOMENA OF CAST HIGH ALLOY HEAT RESISTANT ALLOY HK-40 AND CORROSION RESISTANT 19CR-9NI2.5MO STEELS," REPORT A-69, STEEL FOUNDERS' SOCIETY OF AMERICA, 1980 9. C.H. CADDEN, "WELDABILITY OF HIGH ALLOY CASTINGS," M.S. THESIS, RENSSELAER POLYTECHNIC INSTITUTE, TROY, NEW YORK, MAY 1981 10. D. ROZET, H.C. CAMPBELL, AND R.D. THOMAS, WELD. J., OCT 1949, P 481-S TO 491-S 11. Y. ARATA, F. MATSUDA, AND S. KATAYAMA, TRANS. JWRI, JUNE 1977, P 105-116 12. M.J. CIESLAK AND W.F. SAVAGE, WELD. J., MAY 1985, P 119-S TO 126-S

Dissimilar Welds With Stainless Steels Joel G. Feldstein, Foster Wheeler Energy Corporation

Introduction THE WELDING OF DISSIMILAR AUSTENITIC STAINLESS STEELS to each other is fairly common practice. When suitable welding procedures and filler metals are employed, most austenitic stainless steels can also be welded satisfactorily to several other classes of weldable steel, including ferritic and precipitation-hardening stainless steels, carbon steels, and low-alloy steels. As a general rule, the weld metal should be at least equal in strength and corrosion properties to the poorest component in the joint. Additionally, there should be no intermetallic compounds or other phases to degrade the properties of the weld metal. Further, the microstructure must be resistant to cracking. An austenitic microstructure with a Ferrite Number (FN) of 5 to 10 has been shown to be highly resistant to cracking. If a fully austenitic weld metal occurs, satisfactory resistance to cracking can be achieved in certain applications by using a filler metal of high manganese content or a nickel-base alloy. Dissimilar Welds With Stainless Steels Joel G. Feldstein, Foster Wheeler Energy Corporation

Welding Dissimilar Austenitic Stainless Steels In general, the deposited weld metal composition should nearly match the base metal composition when welding different austenitic stainless steels to themselves. However, it is not uncommon to weld different austenitic stainless steels, such as types 304 and 347, to each other or to a nickel-base alloy using a filler metal that differs in chemical composition from both base metals provided the weld metal will have suitable corrosion resistance or mechanical properties (or both) for the required service. Generally, the alternate filler metal should be more highly alloyed than the base metal, but the ferrite content of the resulting weld metal should be considered for corrosion or high-temperature service. Some filler metals may give too high a ferrite content, particularly the high-chromium, low-nickel types, such as type 312. Recommended filler metals for welding dissimilar austenitic stainless steels are given in Table 1. Additional information on stainless steel filler metals can be found in the article "Welding of Stainless Steels" in this Volume.

TABLE 1 STAINLESS STEEL FILLER METALS FOR WELDING SIMILAR AND DISSIMILAR AUSTENITIC STAINLESS STEELS

BASE METALS

201,202, 301,302, 302B, 303,304, 305,308 340L 309, 309S 310, 310S, 314 316 316L 317 317L 321,347,348 330 (A) (B)

FILLER METALS 201, 202, 301, 302, 302B, 303, 304, 305, 308

304L

E308(A)

E308(A) E308(A) E308L

309 309S

310 310S, 314 E308(A)

E308(A) E308(A) E309 E309 310

FOR SERVICE TEMPERATURE OF 494

~77-83 >72

460-503 >416

~67-73 >60

~27 27-28

PURE

Ref 42 Resistance welding is a process in which heat is produced by resistance to the flow of high electrical currents (5 to 10 kA) across the interface between two contacting surfaces. At sufficiently high current levels, melting is induced at the interface and a weld fusion zone (i.e., weld nugget) is created. At lower current levels, welding of the faying surfaces occurs entirely by a solid-state deformation or diffusion welding mechanism. Resistance welding is most commonly used to join sheet and can produce either discrete, individual spot welds or continuous, overlapping seam welds. The resistance welding parameters employed for titanium are generally comparable to those used for austenitic stainless steels, due to similarities in thermal conductivity between the two materials. As expected, the production of high-quality resistance welds requires careful cleaning of the sheet surfaces prior to welding. However, the rapid thermal cycles employed and the self-contained nature of the weld zone generally preclude the necessity for inert-gas shielding during welding. Table 12 presents typical parameters for resistance spot welding of Ti-6Al-4V.

TABLE 12 TYPICAL OPERATING CONDITIONS FOR RESISTANCE SPOT WELDING OF TI-6AL-4V

MATERIAL THICKNESS mm in.

WELDING CURRENT, KA

ELECTRODE FORCE(A) kn lbf

0.89 1.57 2.36

5.5 10.6 12.5

2.7 6.7 10.7

0.035 0.062 0.093

600 1500 2400

NUGGET WELD DIAMETER TIME, CYCLES (60 mm in. HZ) 7 6.35 0.25 10 9.15 0.36 16 10.9 0.43

Source: Ref 4

(A) RWMA CLASS 2 COPPER ALLOY ELECTRODES, 75 MM (3 IN.) FACE RADIUS. Diffusion welding is a solid-state joining process that is finding increased application for the fabrication of complex

titanium alloy components. Diffusion welding (or bonding) of titanium involves four basic steps: (1) development of intimate physical contact between the faying surfaces through deformation of asperity peaks (surface roughness) by yielding and creep at low pressures and high temperatures, (2) formation of a metallic bond, (3) diffusion across the faying surface, and (4) grain growth across the original weld interface. Prerequisites for accomplishing these steps include a clean and smooth surface combined with a low applied pressure to promote steps 1 and 2, and moderate-to-high temperatures to promote steps 3 and 4. Titanium alloys are ideal for diffusion welding because titanium exhibits a high solubility for oxygen, which promotes dissolution of surface oxides at elevated temperatures. In addition, the yield and creep strengths of titanium are low at welding temperatures, thereby enhancing stages 1 and 2 of the diffusion welding process. Diffusion welding of Ti-6Al-4V is generally performed at temperatures ranging from about 900 to 950 °C (1650 to 1740 °F), at pressures ranging from 1.3 to 13.8 MPa (200 to 2000 psi), and for times ranging from about 1 to 6 h. In combination with superplastic forming, diffusion welding has been utilized to produce a wide variety of complex Ti-6Al4V shapes, such as honeycomb panels for aerospace applications. Other Weld Processes. A wide variety of more specialized welding processes have been used to join titanium,

including flash-butt welding, continuous-drive and inertia friction welding, ultrasonic welding, and explosive welding. In general, these processes have been effective, producing defect-free joints with properties that depend principally on the cooling rates from high temperatures and the characteristics of the resulting transformed-beta microstructure. Special Welding Precautions and Requirements. Production of high-quality welds in titanium and titanium alloys

requires meticulous preweld cleaning of the workpiece and consumables and adequate gas shielding to prevent atmospheric contamination. Mill-annealed titanium sheet and plate are generally supplied in a condition that requires only

the removal of surface grease, oil, and dirt. Such cleaning is normally accomplished by wiping with or dipping in a nonchlorinated solvent such as acetone or methyl-ethyl-ketone (MEK). Materials that exhibit a light oxide scale as a result of heat treatment below about 600 °C (1110 °F) or that contain entrapped oil from machining operations should be pickled for 5 to 10 min in a solution of 30 to 40% nitric acid and 4 to 5% hydrofluoric acid in water at a temperature between 20 and 70 °C (68 and 160 °F). Heavier oxide scales resulting from high-temperature heat treatment or thermomechanical processing may require cleaning in a molten salt bath or mechanical removal, followed by pickling. Pickled pieces should be rinsed in water, dried, and either directly welded or wrapped in clean plastic. All handling should be performed with clean white gloves. Weld contamination can also result from filler metal that has become dirty due to inadequate storage or careless handling prior to or during the welding operation. Degreasing the filler rod or wire is generally sufficient to eliminate contamination problems originating from these sources. Effective shielding of the weld zone from the atmosphere during welding is extremely important in ensuring maximum weld ductility and toughness and in reducing the potential for solid-state weld cracking. Optimum shielding conditions are provided by welding within a permanent or collapsible clear plastic enclosure that has been evacuated and purged with argon (dew point, -24 °C, or -11 °F). Alternatively, localized trailing and backing shields can be used, particularly for automatic welding of simple joint geometries (for example, automatic butt welding of sheet). Use of localized shielding for manual welding of complex parts, although feasible, requires specially designed shielding fixtures and generally is not recommended. Fortunately, evidence of titanium weld contamination is readily apparent by discoloration of the weld surface, with a change from bright silver to straw, to magenta, to blue, with a gradual increase in the level of contamination. Severe contamination is associated with a white or gray powdery-appearing weld surface, and is often accompanied by solid-state cracking in the weld metal. In many high-performance applications, a bright silver color is required, although a light straw color may be accepted in less-critical applications. Because titanium alloy welds are commonly used in fatigue-critical applications, a stress relief operation is generally required following welding. Specific stress-relief temperatures and times depend on the base-metal microstructure (Table 3). In order to prevent property degradation, stress relieving should be performed on components free of surface contamination and in an inert or vacuum environment. Weld Design Criteria and Limitations. Designers of titanium welded structures must consider both welding process applicability and the physical characteristics and mechanical properties of the welds. From a welding process stand-point, an efficient design will utilize a welding process optimally suited to a particular material thickness and joint configuration. Such process suitability is also based on component size and shape (for example, whether the component will fit in an available EBW chamber, or whether the entire part can be suitably protected from the atmosphere during GTAW or PAW), the cost of producing the weld (including both capital equipment and operating costs), and postweld processing requirements.

The design of titanium welded structures is also influenced by the characteristics and mechanical properties of a joint produced by a particular welding process. Physical characteristics of the weld, including undercut or underfill in fusion welds, upset in solid-state welds, and weld distortion are important considerations, because these characteristics affect not only the physical dimensions of the component but also joint mechanical properties. Obviously, the mechanical properties of the joint, as influenced by the integrity and metallurgical structure of the weld zone, are the principal considerations in weld design. As indicated above, low fusion zone ductility and toughness in rapidly cooled welds, and poor axial fatigue behavior in welds containing defects, are important characteristics of welds that may influence the design of the welded component. Welding Specifications. The American Welding Society has published two specifications related to welding of

titanium. AWS Specification 5.16, "Specification for Titanium and Titanium Alloy Bare Welding Rods and Electrodes," provides requirements for titanium and titanium-alloy welding rods for use with GTAW and GMAW processes (see the article "Welding of Titanium Alloys" in this Volume for filler metal compositions). Although the basic specification principally covers filler metal chemical composition requirements, a rather lengthy appendix presents additional general information on welding of titanium, including toughness data for alpha-beta titanium alloy welds. AWS Specification D10.6, "Recommended Practices for Gas Tungsten-Arc Welding of Titanium Pipe and Tubing," is specific to singlephase alpha alloys. Numerous additional specifications developed by the government and by industry discuss the procedural aspects of the welding of titanium and set forth minimum weld quality and mechanical property requirements. These specifications may be quite general or very specific, describing in detail the fabrication or weld repair of an individual component.


4. TITANIUM, ZIRCONIUM, HAFNIUM, TANTALUM, AND COLUMBIUM, WELDING HANDBOOK, 7TH ED., VOL 4, AMERICAN WELDING SOCIETY, 1982, P 434-470 41. R.F. VAUGHAN, PROPERTIES OF WELDED TITANIUM ALLOYS AND THEIR APPLICATION IN THE AEROSPACE INDUSTRY, TITANIUM '80 SCIENCE AND TECHNOLOGY, VOL 4, H. KIMURA AND O. IZUMI, ED., THE METALLURGICAL SOCIETY OF AIME, 1983, P 2423 42. J. MAZUMDER AND W.M. STEEN, MICROSTRUCTURE AND MECHANICAL PROPERTIES OF LASER WELDED TI-6AL-4V, METALL. TRANS. A, VOL 13, 1982, P 865-871 Selection and Weldability of Conventional Titanium Alloys W.A. Baeslack III, The Ohio State University; J.R. Davis, Davis and Associates; C.E. Cross, Martin Marietta Astronautics Group

References

1. M.J. DONACHIE, JR., TITANIUM, METALS HANDBOOK DESK EDITION, AMERICAN SOCIETY FOR METALS, 1985, P 9-1 TO 9-3 2. S. LAMPMAN, WROUGHT TITANIUM AND TITANIUM ALLOYS, PROPERTIES AND SELECTION: NONFERROUS ALLOYS AND SPECIAL-PURPOSE MATERIALS, VOL 2, ASM HANDBOOK, ASM INTERNATIONAL, 1990, P 592-633 3. E.W. COLLINGS, THE PHYSICAL METALLURGY OF TITANIUM ALLOYS, AMERICAN SOCIETY FOR METALS, 1984 4. TITANIUM, ZIRCONIUM, HAFNIUM, TANTALUM, AND COLUMBIUM, WELDING HANDBOOK, 7TH ED., VOL 4, AMERICAN WELDING SOCIETY, 1982, P 434-470 5. W.A. BAESLACK III, D.W. BECKER, AND F.H. FROES, ADVANCES IN TITANIUM ALLOY WELDING METALLURGY, J. MET., MAY 1984, P 46-58 6. R.W. SCHUTZ AND D.E. THOMAS, CORROSION OF TITANIUM AND TITANIUM ALLOYS, CORROSION, VOL 13, ASM HANDBOOK, ASM INTERNATIONAL, 1987, P 669-706 7. D.W. BECKER, W.A. BAESLACK III, AND R.W. MESSLER, JR., TITANIUM WELDING, TITANIUM '80 SCIENCE AND TECHNOLOGY, VOL 1, H. KIMURA AND O. IZUMI, ED., THE METALLURGICAL SOCIETY OF AIME, 1980, P 255-275 8. D.F. HASSON, "A REVIEW OF TITANIUM WELDING PROCESSES," U.S. NAVY REPORT, ANNAPOLIS, MD, 1985 9. J.R. HARTMAN AND W.A. BAESLACK III, ADVANCES IN ARC WELDING OF TITANIUM ALLOYS, WELDABILITY OF MATERIALS, R.A. PATTERSON AND K.W. MAHIN, ED., ASM INTERNATIONAL, 1990, P 319-330 10. W.A. BAESLACK III AND D.L. HALLUM, NATURE OF GRAIN REFINEMENT IN TITANIUM ALLOY WELDS BY MICROCOOLER INOCULATION, WELD. J. RES. SUPPL., VOL 69 (NO. 12), DEC 1990, P 326S-336S 11. D.W. BECKER AND W.A. BAESLACK III, PROPERTY-MICROSTRUCTURE RELATIONSHIPS IN METASTABLE-BETA TITANIUM ALLOY WELDMENTS, WELD. J. RES. SUPPL., VOL 59, 1980, P 85S-92S 12. W.A. BAESLACK III, P.S. LIU, D.P. BARBIS, J.R. SCHLEY, AND J.R. WOOD, POSTWELD HEAT TREATMENT OF GTA WELDS IN A HIGH-STRENGTH, METASTABLE-BETA TITANIUM ALLOY-BETA-C, PROC. OF 1992 INT. CONF. ON TITANIUM ALLOYS (SAN DIEGO), 1992 13. A.T. D'ANNESSA, REDISTRIBUTION OF SOLUTE IN FUSION WELDING, WELD. J., VOL 45, 1966, P 569S-576S 14. W.A. BAESLACK III, EFFECT OF SOLUTE BANDING ON TRANSFORMATIONS IN DISSIMILAR

TITANIUM ALLOY WELDMENTS, J. MAT. SCI. LET., VOL 1, 1982, P 229-231 15. P.S. LIU, W.A. BAESLACK III, AND J. HURLEY, DISSIMILAR ALLOY LASER WELDING OF TITANIUM: TI-6AL-4V TO BETA-C, WELD. J. RES. SUPPL., IN PRESS 16. P.S. LIU, W.A. BAESLACK III, AND J. HURLEY, LASER WELDING OF TI-6AL-4V TO BETA-21S, PROC. OF TMS CONF. ON BETA TITANIUM ALLOYS (DENVER, CO), TMS, 1993 17. W.A. BAESLACK III, S. CHIANG, AND C.A. ALBRIGHT, J. MAT. SCI. LET, VOL 9, 1990, P 698-702 18. W.A. BAESLACK ILL, METALLOGRAPHY, VOL 13, 1980, P 277-281 19. N.N. PROKHOROV, WELD. PROD., VOL 9, 1986, P 30-32 20. P.W. HOLSBERG, "CRACKING AND POROSITY IN TITANIUM ALLOY WELDMENTS," U.S. NAVY DTNSRDC REPORT, 1979 21. F.D. ROSI, D.A. DUBE, AND B.H. ALEXANDER, TRANS. AIME, VOL 197, 1953, P 257 22. N.E. PATON AND J.C. WILLIAMS, THE EFFECT OF HYDROGEN ON TITANIUM AND ITS ALLOYS, HYDROGEN IN METALS, AMERICAN SOCIETY FOR METALS, 1974, P 409-431 23. R.E. LEWIS, I.L. CAPLAN, AND W.C. COONS, THE ELEVATED TEMPERATURE DUCTILITY DIP PHENOMENON IN ALPHA, NEAR-ALPHA AND ALPHA-BETA TITANIUM ALLOYS, TITANIUM SCIENCE AND TECHNOLOGY, VOL 2, G. LÜTJERING, U. ZWICKER, AND W. BUNK, ED., DEUTSCHE GESELLSCHAFT FÜR METALLKUNDE E.V., WEST GERMANY, 1985 24. B.K. DAMKROGER, G.R. EDWARDS, AND B.B. RATH, A MODEL FOR HIGH TEMPERATURE DUCTILITY LOSS IN ALPHA-BETA TITANIUM ALLOYS, METALL. TRANS. A, VOL 18, 1987, P 483-485 25. D.M. BOWDEN AND E.A. STARKE, JR., HOT CRACKING SUSCEPTIBILITY OF TITANIUM 6211, TITANIUM SCIENCE AND TECHNOLOGY, VOL 2, G. LÜTJERING, U. ZWICKER, AND W. BUNK, ED., DEUTSCHE GESELLSCHAFT FÜR METALLKUNDE E.V., WEST GERMANY, 1985, P 853 26. D.M. BOWDEN AND E.A. STARKE, JR., METALL. TRANS. A, VOL 15, 1984, P 1687-1698 27. D. HAYDUK, B.K. DAMKROGER, G.R. EDWARDS, AND D.L. OLSON, WELD. J., VOL 65, 1986, P 251S-260S 28. B.K. DAMKROGER, G.R. EDWARDS, AND B.B. RATH, WELD. J., VOL 68, 1989, P 290S-302S 29. R.E. TREVISAN, D.D. SCHWEMMER, AND D.L. OLSON, THE FUNDAMENTALS OF WELD PORE FORMATION, WELDING: THEORY AND PRACTICE, ELSEVIER, 1990, P 79-115 30. K. BORGGREEN AND I. WILSON, USE OF POSTWELD HEAT TREATMENTS TO IMPROVE DUCTILITY IN THIN SHEETS OF TI-6AL-4V, WELD. J., VOL 59, 1980, P 1S-8S 31. S.L. BOSTON AND W.A. BAESLACK III, "MICROSTRUCTURE/MECHANICAL PROPERTY RELATIONSHIPS IN GTA WELDED TI-10V-2FE-3AL," AFML-TM-LL-80-1, WRIGHT PATTERSON AIR FORCE BASE, DAYTON, OH, 1980 32. M.A. GREENFIELD AND C.M. PIERCE, POSTWELD AGING OF A METASTABLE BETA TITANIUM ALLOY, WELD. J., VOL 52, 1973, P 524S-527S 33. W.A. BAESLACK III, EVALUATION OF TRIPLEX HEAT TREATMENTS FOR TITANIUM ALLOYS, WELD. J., VOL 61, 1982, P 193S-199S 34. D.W. BECKER AND W.A. BAESLACK III, PROPERTY/MICROSTRUCTURE RELATIONSHIPS IN METASTABLE-BETA TITANIUM ALLOY WELDMENTS, WELD. J., VOL 59, 1980, P 850-920 35. W.A. BAESLACK III AND F.D. MULLINS, RELATIONSHIP OF MICROSTRUCTURE, MECHANICAL PROPERTIES AND FRACTOGRAPHIC FEATURES IN WELDED HIGH TOUGHNESS TITANIUM ALLOYS, TRENDS IN WELDING RESEARCH, S.A. DAVID, ED., AMERICAN SOCIETY FOR METALS, 1982, P 541 36. V.N. ZAMKOV ET AL., WAYS OF PROVIDING THE HIGH SERVICEABILITY OF HEATHARDENED WELDED JOINTS OF ALPHA + BETA TITANIUM ALLOYS, TITANIUM SCIENCE AND TECHNOLOGY, VOL 2, G. LÜTJERING, U. ZWICKER, AND W. BUNK, ED., DEUTSCHE GESELLSCHAFT FÜR METALLKUNDE, E.V., WEST GERMANY, 1985, P 877 37. W.A. BAESLACK III, D.W. BECKER, AND F.H. FROES, WELDABILITY OF A HIGH TOUGHNESS

TITANIUM ALLOY, A METALLURGICAL EVALUATION, TITANIUM '80 SCIENCE AND TECHNOLOGY, VOL 2, H. KIMURA AND O. IZUMI, ED., THE METALLURGICAL SOCIETY OF AIME, 1980, P 2369 38. R.E. LEWIS ET AL., THE EFFECT OF BORON ON WELDMENT MICROSTRUCTURES IN TI-6AL2NB-1TA-1MO ALLOY, TITANIUM SCIENCE AND TECHNOLOGY, VOL 2, G. LÜTJERING, U. ZWICKER, AND W. BUNK, ED., DEUTSCHE GESELLSCHAFT FÜR METALLKUNDE E.V., WEST GERMANY, 1985, P 799 39. T.S. PARKER AND P.G. PARTRRIDGE, FATIGUE PROPERTIES OF 4 MM THICK LASER WELDED TI-6AL-4V SHEET, FATIGUE '84, C.J. BEEVERS, ED., ENGINEERING MATERIALS ADVISORY SERVICES LTD., LONDON, 1984 40. P.S. PAO, T.C. PENG, AND J.E. O'NEAL, FATIGUE CRACK PROPAGATION THROUGH TITANIUM WELDS, TITANIUM SCIENCE AND TECHNOLOGY, VOL 2, G. LÜTJERING, U. ZWICKER, AND W. BUNK, ED., DEUTSCHE GESELLSCHAFT FÜR METALLKUNDE, E.V., WEST GERMANY, 1985, P 887 41. R.F. VAUGHAN, PROPERTIES OF WELDED TITANIUM ALLOYS AND THEIR APPLICATION IN THE AEROSPACE INDUSTRY, TITANIUM '80 SCIENCE AND TECHNOLOGY, VOL 4, H. KIMURA AND O. IZUMI, ED., THE METALLURGICAL SOCIETY OF AIME, 1983, P 2423 42. J. MAZUMDER AND W.M. STEEN, MICROSTRUCTURE AND MECHANICAL PROPERTIES OF LASER WELDED TI-6AL-4V, METALL. TRANS. A, VOL 13, 1982, P 865-871 Selection and Weldability of Advanced Titanium-Base Alloys David H. Phillips, Edison Welding Institute

Introduction ADVANCED TITANIUM-BASE ALLOYS can generally be categorized as having conventional (terminal phase), intermetallic, or titanium-matrix composite structures. Conventional alloys can be further subdivided, on the basis of the amount of beta phase that is present, into four families: near-alpha, alpha-beta, near-beta, or metastable-beta (Table 1). Intermetallics that are based on titanium most commonly include alpha-2, gamma, and orthorhombic systems. Although the phrase titanium-matrix composite covers a potentially wide range of matrix materials and either continuous or discontinuous reinforcements, the most common composite used is continuous silicon-carbide-reinforced Ti-6Al-4V. This article will describe the physical metallurgy of each alloy category and its families, followed by a brief discussion on the weldability of each.

TABLE 1 ADVANCED TITANIUM ALLOY FAMILIES

CONVENTIONAL ALLOYS NEAR-ALPHA: TI-1100; IMI-834 ALPHA-BETA: CORONA 5 NEAR-BETA: TI-10V-2FE-3AL METASTABLE-BETA: BETA 215; BETA C INTERMETALLIC ALLOYS ALPHA-2: TI24AL-11NB; TI25AL-10NB-3V-1MO GAMMA: TI-AL-NB-CR ORTHORHOMBIC: TI-AL-NB TITANIUM-MATRIX COMPOSITES TI-6AL-4V WITH SIC CONTINUOUS FIBERS

Selection and Weldability of Advanced Titanium-Base Alloys David H. Phillips, Edison Welding Institute

Physical Metallurgy Conventional alloys can usually be described by one of the four families shown in Table 1. The compositions and maximum use temperatures of selected alloys from these families are shown in Table 2, whereas basic mechanical properties are shown in Table 3.

TABLE 2 MAXIMUM USEFUL TEMPERATURES AND CHEMICAL COMPOSITIONS OF SELECTED HIGHTEMPERATURE TITANIUM ALLOYS

ALLOY DESIGNATION MAXIMUM USEFUL TEMPERATURE °C °F TI-1100 595 1100 IMI-834 595 1100 CORONA 5 ... ... TI-10V-2FE-3AL 315 600 BETA 21S ... ... BETA C ... ...

COMPOSITION, WT% Al 6 5.5 4.5 3 3 3

Sn 2.75 4 ... ... ... ...

Zr 4 4 ... ... ... 4

Mo 0.4 0.3 5 ... 15 4

Nb ... 1 ... ... 2.7 ...

V ... ... ... 10 ... 8

Si 0.45 0.5 ... ... 0.2 ...

Cr ... ... 1.5 ... ... 6

C ... 0.06 ... ... ... ...

TABLE 3 BASIC MECHANICAL PROPERTIES OF ADVANCED TITANIUM-BASE ALLOYS COMMON NAME

TI-1100 IMI 834 CORONA 5 TI-10V-2FE-3AL BETA 21S BETA C ALPHA-2 GAMMA ORTHORHOMBIC

ULTIMATE TENSILE STRENGTH MPa Ksi

MPa

Ksi

1010 1050 1200 1000 1100 1300 690 480 1100

910 950 1170 1000 1000 1200 590 340 1030

130 140 170 140 140 175 85 50 150

145 150 175 145 160 190 100 70 160

YIELD STRENGTH


ELONGATION, %

FRACTURE TOUGHNESS MPa m

Ksi in

21 24 22 50 ... ... 4 2 5

... ... ... ... 9 9 ... ... ...

67 44 ... 55 ... 88 22 16 22

61 40 ... 50 ... 80 20 15 20

The physical metallurgies of selected alloys from each family are described below. Near-Alpha Alloys. The alloy Ti-1100 was developed in response to a requirement for higher creep properties in Ti6242S. It was determined that, by adjusting the tin and molybdenum compositions and adding silicon, significant improvements to the creep properties could be realized (Ref 1). Because it was developed for a maximum use temperature of 595 °C (1100 °F), it was named Ti-1100. The processing of this alloy usually involves a final beta forge, followed by an 8-h stabilization at 600 °C (1110 °F). This results in a fully transformed acicular structure, which is the proper microstructure for most applications.

The near-alpha alloy IMI-834 was developed for applications requiring a better balance of creep and fatigue properties. Because fully transformed microstructures are not ideal, in terms of fatigue, IMI-834 was developed to facilitate processing into a duplex microstructure. This was achieved by adding carbon, which stabilizes the alpha phase, thus

increasing the size of the alpha-beta phase field. As a result, although careful control is required, a duplex microstructure can be achieved through subsequent alpha-beta processing (Ref 2). The microstructure typically contains 15% primary alpha. Niobium was added to increase oxidation resistance during high-temperature exposures, and silicon was added to improve strength and creep resistance. Alpha-Beta Alloys. The need for titanium alloys with higher levels of toughness led to the development of Corona 5. This alpha-beta titanium alloy has intermediate strength and exceptionally high fracture toughness. In order to achieve the latter property, this alloy is typically processed high in the alpha-beta phase field, which results in an alpha and beta microstructure containing moderately high aspect ratio primary alpha particles. One of the principal attractions of this alloy is that its strength level can easily be adjusted by modifying the dispersion of the fine alpha (transformed beta) phase (Ref 3). Near-Beta Alloys. Ti-10V-2Fe-3Al, which is considered to be a relatively high-strength near-beta alloy, can be

processed for either high strength and medium toughness or moderate strength and high toughness. It is typically solution treated and aged, which results in a mostly transformed alpha and beta microstructure. Metastable-Beta Alloys. The Beta 21S alloy was developed primarily as a material that had acceptable high-

temperature properties and oxidation resistance, could be readily cold rolled, and was age hardenable. These features (primarily the cold rollability) make it an excellent candidate as a matrix material for titanium-matrix composites. Basemetal processing usually involves the hot and cold working of the ingot to form billet or strip product, followed by solution annealing, cold forming to near-net shape, resolution annealing, and then aging. The fine alpha platelets that precipitate during the aging process provide increased strength in the alloy. The resultant microstructure contains the fine alpha platelets in a beta-phase matrix. Beta C is a metastable-beta alloy that can be processed and heat treated to provide an excellent combination of yield and tensile strengths, ductility, fracture toughness, and fatigue strength (Ref 4). The alloy contains the beta-isomorphous elements molybdenum and vanadium, as well as the sluggish beta-eutectoid element chromium at levels that depress the beta transus temperature to approximately 750 °C (1380 °F) and promote retention of the high-temperature beta phase down to room temperature upon air cooling. A solution heat treatment, followed by aging, promotes the precipitation of a fine distribution of alpha phase in the beta-phase matrix. Intermetallics. Titanium intermetallics consist primarily of the titanium aluminides alpha-2 and gamma, as well as the

orthorhombic plus beta types of alloys. The aluminides, as shown on the titanium-aluminum phase diagram of Fig. 1, consist of two alloy families: alpha-2, which is based on the Ti3Al intermetallic, and gamma, which is based on the titanium-aluminum intermetallic. These alloys offer greater stiffness and improved specific creep resistance, and they are capable of extending the use temperature range of conventional titanium alloys. The orthorhombic alloys contain large amounts of niobium, which promotes formation of the orthorhombic phase, Ti2AlNb.

FIG. 1 TITANIUM-ALUMINUM PHASE DIAGRAM. SOURCE: REF 5

Alpha-2 alloys represent a growing family of intermetallics that contain from 14 to 25 at.% aluminum, which places

them in the Ti3Al phase field of Fig. 1. They usually contain beta-stabilizing elements, such as niobium and molybdenum, to promote beta-phase retention and ductility. These alloys are typically alpha-beta processed, resulting in a microstructure similar to that of conventional alloys, which consists of islands of equiaxed primary alpha-2 in a matrix of transformed beta. When compared with conventional titanium alloys, alpha-2 alloys offer higher temperature capabilities (from 100 to 150 °C, or 180 to 270 °F, higher), but have reduced room-temperature toughness and ductility properties (3 to 6% elongation). The gamma alloy family, which is growing even more rapidly than the alpha-2 alloy family, typically contains from 45

to 52 at.% aluminum, which places these alloys primarily in the gamma phase field, with some overlap in the gamma/alpha-2 phase field. This results in a range of alloy microstructures, including single-phase gamma (TiAl) and duplex structures containing varying amounts of alpha-2 phase (Ti3Al). Duplex microstructures are generally preferred, because of their higher ductility. Alloying elements such as niobium, chromium, vanadium, tantalum, manganese, and molybdenum are added in small amounts (generally from 2 to 4 at.%) to improve properties such as room-temperature ductility and oxidation resistance. The gamma alloys are more difficult to thermomechanically process than are the alpha2 alloys, primarily because of their very high strengths at elevated temperatures. Accordingly, compared to alpha-2 alloys, the gamma alloys offer temperature capabilities that are from 100 to 150 °C (180 to 270 °F) higher. They also are known to suffer from low fracture toughness and ductility, although aggressive alloy development programs are currently attempting to improve these properties. Orthorhombic alloys represent a new class of alloys based on the orthorhombic and beta phases. These alloys offer

temperature capabilities similar to those of the alpha-2 alloys, but with higher ductilities and toughness at low temperatures. Although these alloys are similar to alpha-2 alloys, a higher niobium content results in an orthorhombic extension of the hexagonal lattice. The large amounts of niobium also result in a highly stabilized beta phase, compared with alpha-2 alloys. Processing is similar to that used with other titanium alloys. Forging is typically conducted below the beta transus temperature, resulting in a two-phase orthorhombic plus transformed-beta microstructure. A direct age then can be subsequently applied to develop properties.

Titanium-Matrix Composites. Although many titanium matrix/fiber combinations are currently being evaluated by

aerospace companies and titanium producers, the most popular composite consists of a conventional Ti-6Al-4V or Ti6242 matrix with layers of continuous silicon-carbide fibers. The layers can be either unidirectional or cross-ply. Two manufacturing techniques can be used: a foil/fiber/foil approach, in which layers of monolithic foil and fiber sets are either vacuum hot pressed or hot isostatically pressed to form the composite, or a plasma-spray approach, in which the composite is built up by sequentially winding individual layers of fibers and then spraying the monolithic layers, followed by consolidation using either vacuum hot pressing or hot isostatic pressing.


1. P. BANIA, NEXT GENERATION TITANIUM ALLOYS FOR ELEVATED TEMPERATURE SERVICE, ISIJ INT., VOL 31, 1991, P 840-847 2. OPTIMIZATION OF CREEP AND FATIGUE RESISTANCE IN HIGH TEMPERATURE TI ALLOYS IMI 829 AND IMI 834, PROC. 5TH INT. CONFERENCE ON TITANIUM (MUNICH, GERMANY), 1984, P 2419 3. D.W. BECKER, W.A. BAESLACK, AND F.H. FROES, POWDER METALLURGY OF TITANIUM ALLOYS, PROC. CONF. AIME, AIME, 1980 4. W.A. BAESLACK, P.S. LIU, D.B. BARBIS, J.R. SCHLEY, AND J.R. WOOD, POSTWELD HEAT TREATMENT OF GTA WELDS IN A HIGH-STRENGTH METASTABLE-BETA TITANIUM ALLOY-BETA-C, INTERNATIONAL TITANIUM CONFERENCE (SAN DIEGO), 1992 5. T.B. MASSALSKI, ED., BINARY ALLOY PHASE DIAGRAMS, 2ND ED., VOL 1, ASM INTERNATIONAL, 1990, P 226 Selection and Weldability of Advanced Titanium-Base Alloys David H. Phillips, Edison Welding Institute

Weldability There are many weldability (fusion welding) concerns associated with advanced titanium alloys. These differ appreciably from the weldability concerns that are characteristic of steels, such as hydrogen cracking, or of superalloys, such as strain age or liquation cracking. The weldability problems associated with advanced titanium alloys are related to their microstructural sensitivity to high heat input and rapid cooling rates. These thermal excursions, which are characteristic of many welding processes, can result in excessive beta grain growth and the formation of undesirable solid-state transformation products. The affected microstructures typically suffer from poor ductility and toughness. Such effects can be minimized, however, through proper welding practices, such as those discussed below. Conventional Alloys. Each of the four families previously identified are described below in terms of weldability. The near-alpha alloy Ti-1100 is generally considered to be a very weldable alloy, because of its high ductility.

However, the excellent creep resistance of this alloy also means that high residual stresses will be developed when welding, especially with thick sections. These stresses may have to be relieved by a postweld heat treatment following welding, in order to avoid cracking. The near-alpha alloy IMI 834 is also considered to be weldable, but the weld fusion zone and heat-affected zone (HAZ) may be void in primary alpha phase. These regions may not retain a good balance of creep and fatigue properties, for which IMI 834 is designed. Alpha-Beta Alloys. The Corona 5 high-toughness alloy can be successfully welded, but it requires carefully controlled

cooling rates, as well as postweld heat-treatment conditions. Excessively high cooling rates can result in the formation of orthorhombic martensite, a structure that exhibits inferior strength and fracture toughness. The aging of this martensite at low and intermediate postweld heat-treatment temperatures results in high strength, but low ductility and fracture

toughness. However, through proper cooling and postweld heat treating, a fine alpha-beta fusion zone and HAZ microstructure that exhibits good mechanical properties can be created. Near-Beta Alloys. The Ti-10V-2Fe-3Al alloy is very weldable in the heat-treated condition. Tensile and fracture

toughness values that are near those of the base metal can be achieved with as-welded Ti-10V-2Fe-3Al. Metastable-Beta Alloys. Although metastable-beta titanium alloys, such as Beta 21S, are generally considered to be

weldable, cooling rates should be kept relatively high, and the weld fusion zone size should be minimized to avoid the formation of coarse columnar beta fusion zone grains and large beta HAZ grains, both of which can degrade ductility (Ref 6). If a relatively fine as-welded grain structure can be achieved, then moderate weld strength and high ductility can be achieved. Postweld heat treating to precipitate the alpha phase can significantly increase weld tensile strengths, but at the expense of ductility. Higher postweld aging temperatures may result in the formation of significant amounts of alpha phase along the beta grain boundaries, which changes the fracture mode, but does not necessarily degrade the properties. Like Beta 21S, the Beta C alloy is considered to be highly weldable, but the fusion zone and HAZ beta grain size should be minimized (Ref 4). As-welded structures are characterized by low strength and high ductility. Relatively low postweld aging treatments can significantly increase strength levels, but ductility levels will be reduced, compared with those of the base metal, because of the large beta grain size and fine alpha phase distribution. Figure 2 shows a Beta C gas-tungsten arc weld microstructure, which reveals the large beta grain size. Higher postweld aging cycles coarsen the alpha phase, which reduces the strength but increases the ductility. As with Beta 21S, the higher aging temperatures can result in significant grain boundary alpha formation, which changes the fracture mode, but does not necessarily degrade the properties.

FIG. 2 MICROSTRUCTURE OF AS-WELDED GAS-TUNGSTEN ARC WELD IN BETA C SHEET

The intermetallic alloys are currently in the early stages of development. As a result, only limited weldability studies

have been conducted to date. However, cursory fundamental studies have provided some indication of their weld sensitivities, and are discussed below. Alpha-2 alloys form a similar weld and HAZ to many conventional alpha-beta titanium alloys. However, because of

limited ductility and notch sensitivity, alpha-2 alloys are susceptible to solid-state fusion zone and HAZ cracking. Preheating to reduce residual stress and promote a more-ductile fusion zone and HAZ microstructure may be required in order to eliminate cracking. Weldability can also be improved by minimizing the size of the weld fusion zone and maximizing the cooling rate through the use of high-energy-density welding processes, such as laser-beam or electronbeam welding. Although this has the effect of retaining the metastable-beta phase, which can then be postweld heat treated, cracking may also occur during the heat treatment. Figure 3 shows cracking in the root of an alpha-2 electronbeam weld.

FIG. 3 CRACKING IN ALPHA-2 ELECTRON-BEAM WELD

Gamma Alloys. The gamma fusion zone and weld HAZ microstructures are significantly different from those of other

titanium alloys. Depending on the weld process and metallographic technique used, the HAZ adjacent to the fusion zone can be nearly indistinguishable from the base-metal microstructure. Excessive grain growth does not appear to be an issue with gamma weldments. Because of low ductility and a high modulus, gamma alloys are very susceptible to weld fusion zone and HAZ cracking, resulting in the requirement for a weld preheat or controlled cooling. Orthorhombic alloys are more ductile than the alpha-2 and gamma alloys, and are therefore easier to weld. Because of

high beta-phase stabilities, a large retained-beta-phase HAZ will result from welding. Postweld heat treatments may be required to achieve an optimum combination of properties. Titanium-Matrix Composites. The weldability issues relative to the monolithic titanium alloys also apply to titanium-

matrix composites, with the additional complication of fiber/matrix interaction. Although some fiber damage from the welding process should not significantly degrade the properties, the damage should, of course, be minimized, where possible. Joint design is especially critical with titanium-matrix composites. The selection of a joining process is also critical, where processes such as brazing, solid-state diffusion welding, or transient liquid phase bonding may be required to minimize fiber damage. These processes eliminate or minimize base-metal melting that would damage the fibers.


4. W.A. BAESLACK, P.S. LIU, D.B. BARBIS, J.R. SCHLEY, AND J.R. WOOD, POSTWELD HEAT TREATMENT OF GTA WELDS IN A HIGH-STRENGTH METASTABLE-BETA TITANIUM ALLOY-BETA-C, INTERNATIONAL TITANIUM CONFERENCE (SAN DIEGO), 1992 6. P.S. LIU, K.H. HOU, J. HURLEY, AND W.A. BAESLACK, LASER WELDING OF AN OXIDATION RESISTANT METASTABLE-BETA TITANIUM ALLOY--BETA-215, INTERNATIONAL TITANIUM CONFERENCE (SAN DIEGO), 1992 Selection and Weldability of Advanced Titanium-Base Alloys David H. Phillips, Edison Welding Institute

References

1. P. BANIA, NEXT GENERATION TITANIUM ALLOYS FOR ELEVATED TEMPERATURE SERVICE, ISIJ INT., VOL 31, 1991, P 840-847 2. OPTIMIZATION OF CREEP AND FATIGUE RESISTANCE IN HIGH TEMPERATURE TI ALLOYS IMI 829 AND IMI 834, PROC. 5TH INT. CONFERENCE ON TITANIUM (MUNICH, GERMANY), 1984, P 2419 3. D.W. BECKER, W.A. BAESLACK, AND F.H. FROES, POWDER METALLURGY OF TITANIUM ALLOYS, PROC. CONF. AIME, AIME, 1980 4. W.A. BAESLACK, P.S. LIU, D.B. BARBIS, J.R. SCHLEY, AND J.R. WOOD, POSTWELD HEAT TREATMENT OF GTA WELDS IN A HIGH-STRENGTH METASTABLE-BETA TITANIUM ALLOY-BETA-C, INTERNATIONAL TITANIUM CONFERENCE (SAN DIEGO), 1992 5. T.B. MASSALSKI, ED., BINARY ALLOY PHASE DIAGRAMS, 2ND ED., VOL 1, ASM INTERNATIONAL, 1990, P 226 6. P.S. LIU, K.H. HOU, J. HURLEY, AND W.A. BAESLACK, LASER WELDING OF AN OXIDATION RESISTANT METASTABLE-BETA TITANIUM ALLOY--BETA-215, INTERNATIONAL TITANIUM CONFERENCE (SAN DIEGO), 1992 Selection and Weldability of Heat-Treatable Aluminum Alloys Richard P. Martukanitz, Applied Research Laboratory, Pennsylvania State University

Introduction THE HEAT-TREATABLE ALUMINUM ALLOYS provide good strength and toughness in engineering applications while maintaining the low density and corrosion resistance of aluminum. These attributes allow the heat-treatable alloys to be used in a wide variety of applications, which include the aerospace, transportation, shipbuilding, tankage, piping, and appliance industries. The majority of these alloys are easily welded by the conventional arc welding processes (gasmetal arc and gas-tungsten arc welding), resistance spot and seam welding processes, as well as the high-energy processes (laser-beam and electron-beam welding). However, these alloys also possess certain characteristics inherent to all aluminum alloys that must be considered during welding (Ref 1). Aluminum possesses a tenacious oxide that not only prevents corrosion, but may also impair welding if not minimized. This is achieved by mechanically or chemically removing heavy oxides prior to welding and using cathodic cleaning during direct current electrode positive (DCEP) gas-metal arc welding (GMAW) or alternating current gas-tungsten arc welding (GTAW). The high thermal conductivity of aluminum alloys requires the use of higher heat input for welding when compared to ferrous alloys. This requires greater welding current during arc welding of aluminum alloys. The high coefficient of thermal expansion of aluminum may also result in greater distortion when compared to ferrous alloys. Hence, aluminum is typically arc welded using high currents and high travel speeds. When in the liquid state, aluminum alloys possess a high solubility for hydrogen and very low solubility after solidification. Hence, proper cleaning and shielding must be utilized to ensure the production of welds free of gas porosity. Finally, aluminum alloys, and particularly the heat-treatable alloys, are sensitive to weld cracking. This condition is most easily avoided by proper filler and base alloy selection and adequate filler metal dilution. Anticipation of these characteristics and general knowledge of these materials allow the heat-treatable aluminum alloys to be readily welded.

Reference

1. P.B. DICKERSON AND B. IRVING, WELDING ALUMINUM: IT'S NOT AS DIFFICULT AS IT SOUNDS, WELD. J., VOL 71 (NO. 4), 1992, P 45

Selection and Weldability of Heat-Treatable Aluminum Alloys Richard P. Martukanitz, Applied Research Laboratory, Pennsylvania State University

Description of Heat-Treatable Aluminum Alloys Designation Systems. The mechanical, physical, and chemical properties of aluminum alloys depend upon

composition and microstructure. The addition of selected elements to pure aluminum greatly enhances its properties and usefulness. Because of this, most applications for aluminum utilize alloys having one or more elemental additions. The major alloying additions used with aluminum are copper, manganese, silicon, magnesium, and zinc. The total amount of these elements can constitute up to 10% of the alloy's composition (all percentages given in weight percent unless otherwise noted). Impurity elements are also present, but their total percentage is usually less than 0.15% in aluminum alloys. The two most commonly used methods of increasing the strength of aluminum alloys are to disperse alloying elements in solid solution and cold work the alloy, or to dissolve the alloying elements into solid solution and precipitate them throughout the alloy's matrix as strengthening particles (Ref 2). Alloys in which the first method applies are referred to as the work-hardenable alloys, and the latter method describes the precipitation-hardenable or heat-treatable alloys. The designations used throughout North America for wrought aluminum alloys involve four digits, with the first digit indicating the major alloying additions, such as 6061 alloy representing the 6xxx alloy series. Aluminum casting alloy designations also involve four digits and may be proceeded by a letter to distinguish impurity levels or modifications of an original classification, such as A356.0 casting alloy. It should be pointed out that the fourth digit, which is separated from the others by a decimal point, indicates the product form (ingot or casting). The designations representing the heattreatable wrought aluminum alloys are the 2xxx series (A1-Cu), the 6xxx series (Al-Mg-Si), and the 7xxx series (Al-Zn). The various combinations of alloying additions and strengthening mechanisms used for wrought aluminum alloys are shown in Table 1. The 8xxx aluminum alloy series denotes "other" major alloying constituents and may be found in both the work-hardenable and precipitation-hardenable classes of alloys. Table 2 provides the nominal composition of some of the common heat-treatable aluminum alloys (Ref 3). Included in Table 2 are compositions of selected heat-treatable casting alloys (Ref 4) and the recently commercialized aluminum-lithium alloys (Ref 5). Welding characteristics of aluminum-lithium alloys are described in the article "Selection and Weldability of Aluminum-Lithium Alloys" in this Volume.

TABLE 1 ALLOY DESIGNATIONS OF WROUGHT ALUMINUM ALLOYS

ALLOY SYSTEM ALUMINUM SERIES WORK-HARDENABLE ALLOYS PURE AL 1XXX AL-MN 3XXX AI-SI 4XXX AL-MG 5XXX AL-FE 8XXX AL-FE-NI 8XXX PRECIPITATION-HARDENABLE ALLOYS AL-CU 2XXX AL-CU-MG 2XXX AL-CU-LI 2XXX AL-MG-SI 6XXX AL-ZN 7XXX AL-ZN-MG 7XXX AL-ZN-MG-CU 7XXX AL-LI-CU-MG 8XXX TABLE 2 NOMINAL COMPOSITION OF SELECTED COMMON HEAT-TREATABLE ALUMINUM

ALLOYS

ALLOY COMPOSITION, WT% Si Cu Mn Mg Cr WROUGHT ALLOYS 2014 0.8 4.4 0.8 0.5 . . . 2024 . . . 4.4 0.6 1.5 . . . 2090 . . . 2.7 . . . 0.2 . . . 2095 . . . 4.0 . . . 0.4 . . . 2219 . . . 6.3 0.3 . . . . . . 2519 . . . 5.8 0.3 0.2 . . . 6005 0.8 . . . . . . 0.5 . . . 6009 0.8 0.4 0.5 0.6 0.1 6010 1.0 0.4 0.5 0.8 0.1 6013 0.8 0.9 0.4 1.0 . . . 6061 0.6 0.3 . . . 1.0 0.2 6063 0.4 . . . . . . 0.7 . . . 6262 0.6 0.3 . . . 1.0 0.1 7005 . . . . . . 0.5 1.4 0.1 7039 . . . . . . 0.3 2.8 0.2 7050 . . . 2.3 . . . 2.2 . . . 7075 . . . 1.6 . . . 2.5 0.2 7178 . . . 2.0 . . . 2.8 0.2 8090 . . . 1.0 . . . 1.0 . . . CASTING ALLOYS 242.0 . . . 4.0 . . . 1.5 . . . 319.0 6.0 3.5 . . . . . . . . . 355.0 5.0 1.2 . . . 0.5 . . . 356.0 7.0 . . . . . . 0.4 . . . 357.0 7.0 . . . . . . 0.6 . . .

Zn

OTHER

... ... ... ... 0.06 ... 0.25 0.25 ... ... ... ... 4.5 4.0 6.2 5.6 6.8 ...

... 2.3 LI, 0.12 ZR 1.0 LI, 0.4 AG 0.06 TI, 0.18 ZR 0.06 TI, 0.17 ZR ... ... ... ... ... ... 0.6 PB, 0.6 BI 0.04 TI, 0.14 ZR ... 0.12 ZR ... ... 2.5 LI, 0.10 ZR

... ... ... ... ...

2.0 NI ... ... ... ...

Source: Ref 3, 4, 5

In commercial product forms, such as aluminum sheet, plate, extrusions, and forgings, these alloys have also undergone varying degrees of mechanical working at ambient and elevated temperatures. It is the combination of alloying additions and thermal-mechanical processing that imparts the desired characteristics to the product. Common considerations in designing alloy compositions and processing practices are strength, toughness, and resistance to corrosion. The temper designation, which follows the alloy designation, is used to describe the thermal and mechanical processing history of the material. The common temper designations for aluminum alloys utilize an H to denote work-hardening, a T to describe various heat-treatment sequences, or an O to represent annealed material. The most common product forms of the heat-treatable alloys dictate the use of either the T or O tempers, such as 6061-T6 or 6061-O. One or more numbers may follow the T to indicate the specific sequence of thermal treatments. Table 3 summarizes the numerical designations used to describe the specific heat-treating sequences of the T-tempers (Ref 6). Additional information on systems for designating aluminum and aluminum alloys can be found in the article "Alloy and Temper Designation Systems for Aluminum and Aluminum Alloys" in Volume 2 of the ASM Handbook.

TABLE 3 NUMERICAL DESIGNATIONS FOR THE HEAT-TREATED TEMPERS OF ALUMINUM ALLOYS

TEMPER DESCRIPTION T1 COOLED FROM AN ELEVATED-TEMPERATURE SHAPING PROCESS AND NATURALLY AGED TO A SUBSTANTIALLY STABLE CONDITION

T2 T3 T4 T5 T6 T7 T8 T9 T10

COOLED FROM AN ELEVATED-TEMPERATURE SHAPING PROCESS, COLD WORKED, AND NATURALLY AGED TO A SUBSTANTIALLY STABLE CONDITION SOLUTION HEAT TREATED, COLD WORKED, AND NATURALLY AGED TO A SUBSTANTIALLY STABLE CONDITION SOLUTION HEAT TREATED AND NATURALLY AGED TO A SUBSTANTIALLY STABLE CONDITION COOLED FROM AN ELEVATED-TEMPERATURE SHAPING PROCESS AND THEN ARTIFICIALLY AGED SOLUTION HEAT TREATED AND THEN ARTIFICIALLY AGED SOLUTION HEAT TREATED AND OVERAGED/STABILIZED SOLUTION HEAT TREATED, COLD WORKED, AND THEN ARTIFICIALLY AGED SOLUTION HEAT TREATED, ARTIFICIALLY AGED, AND THEN COLD WORKED COOLED FROM AN ELEVATED-TEMPERATURE SHAPING PROCESS, COLD WORKED, AND THEN ARTIFICIALLY AGED

Source: Ref 6 Metallurgy. Wrought alloys that constitute heat-treatable (precipitation-hardenable) aluminum alloys include the 2xxx,

6xxx, 7xxx, and some of the 8xxx series alloys. They develop their maximum strength from appropriate solution, quenching, and precipitation treatments. The precipitation mechanism requires alloying elements with appreciable solid solubility in aluminum at elevated temperatures, but with limited solubility at lower temperatures. The solution heat treating operation, which involves heating the alloy to an elevated temperature, allows the alloyed elements to form a solid solution with aluminum. The temperature is then quickly reduced by quenching the material in water to create a supersaturated solid solution. A lower-temperature precipitation process, also referred to as artificial aging, is then used to precipitate strengthening phases throughout the metal in a controlled fashion. The driving force for precipitation is the supersaturation present after quenching. This procedure is illustrated using the partial phase diagram for the aluminum-copper system in Fig. 1. Superimposed on the phase diagram are the temperature ranges used for solution heat treating and precipitation heat treating for a composition representing an aluminum alloy containing 4.0% Cu. Increasing the temperature of the 4.0% Cu alloy to 525 °C (980 °F) causes the copper to form a solid solution. When quenched to room temperature, the solid solution becomes supersaturated. If allowed to become stable, this condition is referred to as a T4 temper. Aging the alloy at a temperature of 175 °C (350 °F) causes nucleation and growth of strengthening precipitate. After solution heat treating and aging, the alloy condition is described as being in the T6 temper.

FIG. 1 PARTIAL PHASE DIAGRAM FOR THE ALUMINUM-COPPER SYSTEM SHOWING TEMPERATURE RANGES FOR SOLUTION AND PRECIPITATION HEAT TREATING

Many times these strengthening precipitates are not the equilibrium phases, such as CuAl2 for the case of the 4.0% Cu alloy, but metastable phases. The metastable precipitates that form are determined by the time and temperature of the precipitation heat treating practice. The precipitation sequence for the aluminum-copper system is shown below:

SOLID SOLUTION → GP1 → GP2 + θ'' → θ' → θ (CuAl2)

(EQ 1)

In this system, strengthening is due to the formation of localized concentrations of copper atoms forming Guinier-Preston zones (GP1 and GP2) that are structurally coherent with the aluminum matrix. At longer aging times and higher aging temperatures the number of Guinier-Preston zones increase and lead to increased strength; the maximum strength is regarded as the peakaged condition. At still higher temperatures or longer times the Guinier-Preston zones are replaced by the noncoherent, metastable precipitates (θ'' and θ'). The decrease in the amount of Guinier-Preston zones formed and the loss of particle coherency result in a subsequent reduction of strength, termed the overaged condition. Figure 2 shows the relationship between strength and precipitate formation during aging of an aluminum 4.0% Cu alloy at 190 °C (375 °F) and 130 °C (265 °F) (Ref 7).

FIG. 2 RELATIONSHIP BETWEEN STRENGTH AND PRECIPITATE FORMATION DURING AGING OF AN AL-4CU ALLOY. SOURCE: REF 7

Characteristics and Applications. Alloying additions and subsequent thermal and mechanical processing dictate characteristics such as strength, toughness, and corrosion resistance of the heat-treatable aluminum alloys. However, it is usually the combination of these properties, that is, strength versus toughness, that is controlled, depending on application requirements. The utility of the heat-treatable aluminum alloys is greatly enhanced by the wide variety of product forms available. This includes sheet, plate, rod, bar, extruded shapes, forgings, and castings.

The 2xxx series alloys, having copper as the primary alloying addition, possess high strength but somewhat lower corrosion resistance than most other aluminum alloys. Many of these alloys also possess relatively good elevatedtemperature strength. Magnesium is also added to the 2xxx series alloys for increased strength but results in greater weld crack sensitivity. Alloy 2024 is one of the highest strength 2xxx alloys and is used extensively in the aircraft industry; however, the 1.5% nominal Cu content of 2024 hinders its weldability due to increased weld crack sensitivity. Alloys having controlled levels of magnesium, such as 2014, 2219, and 2519, are utilized for applications requiring good weldability. Generally, these alloys provide good machinability, but less formability, than other heat-treatable aluminum

alloys. Some applications of the 2xxx alloys are aircraft skin sheet, automotive panels, ballistic armor, and forged and machined components. The 6xxx series alloys containing magnesium and silicon provide moderate strengths and good corrosion resistance in relation to other heat-treatable aluminum alloys. Because they are easily extruded, they are available in a wide range of structural shapes, as well as sheet and plate products. The versatility of these alloys is represented in 6061, which is one of the most commonly used aluminum alloys. Typically, the 6xxx alloys have good formability and good weldability. The 6xxx series alloys are used in autobody sheet, structural members, architectural panels, piping, marine applications, screw machine stock, and many other applications. The 7xxx series provides the highest strength of all aluminum alloys, although this claim is being challenged by the recently developed aluminum-lithium alloys. High strength in the 7xxx alloys is achieved by alloying additions of zinc, magnesium, and often copper, combined with controlled thermal and mechanical processing. Copper, in combination with zinc and magnesium in the 7xxx series alloys, increases strength but hampers weldability due to increased susceptibility to weld cracking. Alloy 7075, containing nominally 5.6% Zn, 2.5% Mg, and 1.6% Cu, is a common alloy of this system having a propensity for weld cracking. Alloys 7004, 7005, and 7039 were designed for applications requiring high strength and good weldability. These alloys have limited amounts of copper that reduce the sensitivity to weld cracking. One important characteristic of the 7xxx series is the ability of the heat-affected zone (HAZ) in these alloys to naturally age, or become precipitation-strengthened at room temperature, after welding. Because of this, weld properties continue to improve for up to 30 days after welding. These alloys typically provide good machinability and adequate formability. Applications for the 7xxx alloys include aircraft construction, truck trailers, railcars, and armor plate.


2. ALUMINUM: PROPERTIES AND PHYSICAL METALLURGY, J.E. HATCH, ED., AMERICAN SOCIETY FOR METALS, 1984, P 222 3. ALUMINUM STANDARDS AND DATA, THE ALUMINUM ASSOCIATION, INC., 1988, P 15 4. ALUMINUM: PROPERTIES AND PHYSICAL METALLURGY, J.E. HATCH, ED., AMERICAN SOCIETY FOR METALS, 1984, P 320 5. R.S. JAMES, ALUMINUM-LITHIUM ALLOYS, VOL 2, ASM HANDBOOK, ASM INTERNATIONAL, 1990, P 178-199 6. ALUMINUM STANDARDS AND DATA, THE ALUMINUM ASSOCIATION, INC., 1988, P 11 7. J.M. SILCOCK, T.J. HEAL, AND H.K. HARDY, STRUCTURAL AGEING CHARACTERISTICS OF BINARY ALUMINUM-COPPER ALLOYS, J. INST. MET., VOL 82, 1953-1954, P 239 Selection and Weldability of Heat-Treatable Aluminum Alloys Richard P. Martukanitz, Applied Research Laboratory, Pennsylvania State University

Welding of Heat-Treatable Aluminum Alloys As stated earlier, the heat-treatable aluminum alloys possess characteristics inherent to all aluminum alloys that must be considered before successful welding can be employed. These characteristics include a tenacious oxide, high thermal conductivity, high coefficient of thermal expansion, solidification shrinkage almost twice that of ferrous alloys, relatively wide solidification-temperature ranges, and high solubility of hydrogen when in the molten state. The first two items may be addressed by weld pretreatment and processing conditions--proper weld pretreating to minimize the oxide and welding at high currents and high travel speeds to ensure that adequate penetration is utilized. However, the remaining characteristics, if not carefully considered, may result in the production of welds of unacceptable quality. The high coefficient of thermal expansion and solidification shrinkage, combined with the relatively wide solidificationtemperature range, makes these alloys susceptible to weld cracking. The high solubility of hydrogen in molten aluminum can result in gas porosity unless proper precautions are heeded. In addition, the degree of degradation of the HAZ for these alloys may differ markedly, depending upon weld processing conditions. An understanding of these aspects yields welds in the heat-treatable aluminum alloys having excellent joint properties.

Crack Sensitivity during Welding Weld cracking in aluminum alloys is of concern due to aluminum's relatively high thermal expansion, large change in volume upon solidification, and wide solidification-temperature range. The weld crack sensitivity of heat-treatable aluminum alloys is especially of concern due to the greater amounts of alloying additions used for these alloys. Because of the detrimental effect of weld cracks on joint properties, the weldability of aluminum alloys is defined as its resistance to weld cracking. Weld cracking in aluminum alloys may be classified into two primary categories based on the mechanism responsible for cracking and crack location. Solidification cracking takes place within the weld fusion zone and typically appears along the center of the weld or at termination craters. Liquation cracks occur adjacent to the fusion zone and may or may not be readily apparent. Solidification cracking, or hot tearing, occurs when high levels of thermal stress and solidification shrinkage are present

while the weld pool is undergoing various degrees of solidification. The hot tearing sensitivity of any given aluminum alloy is influenced by a combination of mechanical, thermal, and metallurgical factors (Ref 8). The degree of restraint of welded assemblies plays a significant role in crack sensitivity by increasing the stress on the solidifying weld. Hot tearing occurs within the weld fusion zone and is affected by weld-metal composition and welding parameters. High heat inputs, such as high currents and slow welding speeds, are believed to contribute to weld solidification cracking (Ref 9). It follows that processes that result in minimal heat input, such as electron-beam welding, reduce weld crack sensitivity. The primary method for eliminating cracking in aluminum welds is to control weld-metal composition through filler alloy additions. Crack sensitivity, determined experimentally as a function of weld composition, is shown in Fig. 3 for various binary aluminum systems (Al-Li, Al-Si, Al-Cu, Al-Mg, and the quasibinary Al-Mg2Si) (Ref 10, 11, 12, 13, 14, 15).

FIG. 3 RELATIVE CRACK SENSITIVITY VERSUS WELD COMPOSITION FOR VARIOUS BINARY ALUMINUM SYSTEMS. SOURCE: REF 10, 11, 12, 13, 14, 15

Although most commercial aluminum alloys are not strictly binary systems, crack-sensitivity curves based on binary systems are still useful in selecting a filler alloy to minimize weld crack sensitivity. In complex alloy systems, minor alloying additions may impact crack sensitivity by widening the temperature range in which a coherent solidifying structure is formed. This effect is shown in Fig. 4, which describes the weld crack sensitivity of two commercial quaternary aluminum systems (Al-Zn-Mg-Cu and Al-Mg-Si-Cu) and a ternary system (Al-Cu-Mg) (Ref 16, 17, 18). In all cases, the curves show the dramatic effect of minor alloying additions, such as copper or magnesium, on weld crack sensitivity. Several high-strength, precipitation-hardenable aluminum alloys rely on complex alloying additions to develop properties; however, these additions dramatically increase weld crack sensitivity. The most notable examples are alloys 2024 (Al-4.4Cu-1.5Mg) and 7075 (Al-5.6Zn-2.5Mg-1.6Cu). Increased weld crack sensitivity may also result when welding dissimilar aluminum alloys to produce weld compositions that widen the coherency temperature range, such as welding 2xxx series to 5xxx series alloys (Al-Cu alloys to Al-Mg alloys).

FIG. 4 WELD CRACK SENSITIVITY OF TWO COMMON QUATERNARY ALUMINUM ALLOY SYSTEMS (AL-MG-ZN-CU AND AL-MG-SI-CU) AND A COMMON TERNARY SYSTEM (AL-CU-MG). SOURCE: REF 16, 17, 18

Liquation Cracking. An important element of the HAZ for precipitation-hardenable alloys is the thin boundary layer

adjacent to the fusion zone that is referred to as the partially melted region. This region is produced when eutectic phases or constituents that have low melting points (melting points below the melting point of the bulk material) liquate, or melt, at grain boundaries during welding (Ref 19). It occurs in precipitation-hardenable alloys because of the relatively large amount of alloying additions available to form eutectic phases. During welding, these phases liquate and, if sufficient stress is present, may be accompanied by tears. Under extreme conditions, continuous cracks may form along the fusion zone interface. As expected, higher heat input widens the partially melted region and makes it more prone to cracking. An example of cracks formed in this region for alloy 6013-T6 welded with 5356 filler alloy is shown in Fig. 5 (Ref 20). Filler alloy composition greatly affects whether cracks form in this liquated region (Ref 21 and 22). Filler alloys that have lowsolidification temperatures provide less susceptibility to liquation cracks, since the solidification shrinkage strains occur at lower temperatures. This allows healing, or solidification, of the partially melted region to occur prior to solidification shrinkage strains.

FIG. 5 MICROGRAPH OF A GAS-TUNGSTEN ARC WELD ON 2.0 MM (0.080 IN.) THICK 6013-T6 BASE METAL WITH 5356 FILLER ALLOY SHOWING CRACKS IN THE PARTIALLY MELTED REGION. SOURCE: REF 20

Porosity during Welding Porosity in aluminum weldments occurs when hydrogen gas is entrapped during solidification (Ref 23 and 24). Hydrogen has an appreciable solubility in molten aluminum and a low solubility in the solid. Figure 6 shows the solubility of hydrogen in pure aluminum at temperatures representing the solid and liquid states (Ref 25).

FIG. 6 HYDROGEN SOLUBILITY IN PURE ALUMINUM. SOURCE: REF 25

Hydrogen is absorbed into the molten pool during welding because of its high solubility, and it forms gas pores upon solidification due to the decrease in solubility. It is this difference in solubility that is the driving force for porosity formation (Ref 26). Once a gas bubble forms, it may be expelled from the weld by natural buoyancy or forced convection within the molten pool. Redistribution of hydrogen during solidification may also act to push pores ahead of the solidifying front. Therefore, the welding position and weld parameters may dramatically alter the amount of porosity found in solidified welds. The expulsion of hydrogen pores may be described by the pore's buoyancy velocity and welding speed. If the solidification front is moving at a lower velocity than the pore's buoyancy velocity, gas pores may escape (Ref 27). Lower welding speeds create slower solidification fronts and favor expulsion of gas pores. Welding position may either help or hinder the expulsion of pores. In general, vertical-up welding of aluminum produces the least porosity because the solidifying pool provides easy escape of gas pores. Converse]y, overhead welding produces the greatest amount of porosity because float direction and pool convection patterns cause pores to be entrapped in the weld root during solidification. The sources of hydrogen present in the welding system depend on the particular welding process. In arc welding these sources are: (1) hydrogen from the base metal, (2) hydrogen from the filler metal, and (3) hydrogen within the shielding gas (Ref 28). This is illustrated in Fig. 7 for GMAW. Hydrogen from the base metal may be attributed to the surface in the form of hydrocarbons (lubricants) and hydrated oxides, or it may be internal, as with castings and aluminum powdermetallurgy (PM) parts. Filler metal hydrogen is in the form of hydrocarbons (lubricants) or hydrated oxides on the surface of the filler. Hydrogen in the shielding gas is due to moisture contamination within the gas or aspiration of moist air into the gas system due to defective gas lines and couplings. A dew-point determination of the shielding gas may provide an indication of moisture contamination. Normally, a dew-point measurement of -50 °C (-60 °F) or below, measured at the gas cylinder outlet, is considered sufficient gas quality (Ref 29).

FIG. 7 SOURCES OF HYDROGEN IN GAS-METAL ARC WELDING. HG, HYDROGEN FROM SHIELDING GAS; HE, HYDROGEN FROM ELECTRODE; HB, HYDROGEN FROM BASE METAL. SOURCE: REF 28

In the case of GMAW of aluminum wrought products, the filler electrode, shielding gas, and base-metal surface contaminants may contribute to gas porosity. In the case of welding aluminum castings or PM parts, internal hydrogen of the base metal typically governs pore formation (Ref 30). Because porosity in aluminum welds is attributed to a source of hydrogen contamination within the welding system, preventative measures may be taken to ensure sound welds. This begins with proper precleaning of the parts to be welded. Mill and machining lubricant must be removed by solvent degreasing before welding. Stainless steel wire brushing of the joint to remove thick mill oxides is recommended just prior to welding. During GTAW of aluminum, alternating current with sufficient electrode-positive polarity provides excellent arc cleaning action to remove surface oxides. For GMAW, procurement of quality electrodes and proper electrode storage are necessary to minimize porosity. In all cases, maintenance of gas hoses, regulators, and auxiliary equipment is required to produce porosity-free welds. HAZ Degradation

A HAZ is created adjacent to the fusion zone and results in a degradation of base-metal properties. This degradation is caused by microstructural modifications associated with elevated temperatures experienced in this zone. For heattreatable aluminum alloys, the HAZ is distinguished by dissolution or growth of precipitates. The response in the HAZ of 2xxx aluminum alloys follows dissolution, whereas the primary modification for the 6xxx alloys is growth of precipitates (Ref 31, 32, 33). Although the nature of these HAZs may differ, they are all diffusion-controlled and thermally dependent (Ref 34). A common method of determining the width and extent of the HAZ is by measuring the hardness across this zone. Hardness profiles for gas-tungsten arc welds of two common heat-treatable alloys, 2219-T87 and 6061-T6, are shown in Fig. 8 (Ref 35). A hardness profile for the work-hardenable alloy 5456-H116 is also shown in Fig. 8 for comparison.

FIG. 8 HARDNESS PROFILES ACROSS THE HAZ OF GAS-TUNGSTEN ARC WELDS ON 3.2 MM (0.125 IN.) THICK 2219-T87, 5456-H116, AND 6061-T6 MADE USING CONSTANT HEAT INPUT. SOURCE: REF 35

The hardness profile for 2219-T87 exhibits dissolution of precipitates. The strengthening particles for these alloys are Guinier-Preston zones, which are metastable precipitates. The solvus curves for these phases are superimposed on the partial aluminum-copper phase diagram in Fig. 9 (Ref 31). When the temperatures in the HAZ exceed the various solvi, the respective phases are dissolved. At positions close to the fusion zone, higher temperatures are experienced and greater dissolution of strengthening phases occurs. This results in a continual decrease in strength in the HAZ of aluminumcopper alloys and is illustrated in Fig. 8 for 2219-T87 alloy.

FIG. 9 PARTIAL PHASE DIAGRAM FOR THE ALUMINUM-COPPER SYSTEM SHOWING METASTABLE PHASE SOLVI. SOURCE: REF 31

The hardness profile for the HAZ of 6061-T6 is representative of the 6xxx series alloys and is governed by a growthlike transformation of precipitate. This system also exhibits a precipitation sequence of metastable precipitates, with the principal strengthening phases occurring early in the sequence in the form of Guinier-Preston zones. The transformation of these precipitates at elevated temperatures occurs in a smooth, continuous fashion and is closely akin to growth (Ref 31). The result is the formation of nonstrengthening phases. Due to the competing nature of the diffusion rate and supersaturation with increasing temperatures, the transformation proceeds at the greatest rate when temperatures are between 290 and 425 °C (550 and 800 °F) (Ref 36, 37, 38). The minimum hardness value for 6061-T6 in Fig. 8 represents the position in the HAZ that experiences temperatures within this range. This minimum has also been referred to as the overaged zone for these alloys. At higher temperatures, indicative of positions close to the fusion zone, particles are dissolved into solid solution and upon cooling may precipitate as strengthening phases. This contributes to the slight increase in hardness that typically occurs near the fusion zone for the 6xxx series alloys. Effect of Heat Input on the HAZ. Because the metallurgical transformations of the HAZ of aluminum alloys are

thermally dependent, the welding process and parameters determine the extent of base-metal degradation. High heat input and preheating increase both the degree and width of the HAZ. This is especially true for the precipitation-hardenable aluminum alloys and is illustrated for 6061-T6 welds made using various heat inputs in Fig. 10 (Ref 39). Thus, the HAZ degradation of the heat-treatable alloys can be reduced by employment of multipass welding, close control of interpass temperature, and elimination of preheating. The highest weld strengths for these alloys are achieved when interpass temperatures do not exceed 65 °C (150 °F) and welding heat input is minimized.

FIG. 10 HARDNESS PROFILES OF THE HAZ OF GAS-TUNGSTEN ARC WELDS ON 6061-T6 USING VARIOUS HEAT INPUTS. SOURCE: REF 39

Postweld Heat Treatment and the HAZ Postweld thermal heat treatments can also be used to improve the strength of

the HAZ for heat-treatable alloys. This may involve complete postweld solution heat treating and aging or postweld aging only. Although the recovery of strength in the HAZ after postweld aging is less than postweld solution heat treating and aging, there are advantages to postweld aging only. Temperatures used for postweld aging are much lower than postweld solution heat treating, and postweld aging does not involve water quenching that imposes residual stresses and distortion to welded assemblies. However, great strides have been made in the development and application of polymer quenchants that reduce the amount of distortion observed after quenching and may complement the use of postweld solution heat treating and aging. Figure 11 depicts hardness profiles of the HAZ for 6061-T4 and 6061-T6 starting material in the as-welded and postweld aged conditions (Ref 40). A dramatic recovery of strength in the HAZ is observed when welding 6061-T4 and postweld aging. Postweld thermal treatments may be used to recover strength in the HAZ of the heat-treatable alloys; however, when 6061-T6 is postweld aged, strength in the HAZ is increased while strength in the unaffected base metal is decreased. This is due to overaging of the 6061-T6 base metal; therefore, T4 or T3 material should be utilized when postweld aging the precipitation-hardenable alloys. The 7xxx series alloys possess the ability to naturally age (age at room temperature) in the HAZ. The recovery of strength is similar to postweld artificial aging, but because of the relatively low-temperature aging, the time is substantial. Usually, a period of 14 days is required for improving the HAZ strength of the 7xxx series alloys by natural aging.

FIG. 11 HARDNESS PROFILES OF THE HAZ FOR 6061-T4 AND T6 STARTING MATERIALS IN THE AS-WELDED (AW) AND POSTWELD AGED (PWA) CONDITIONS. SOURCE: REF 40


8. C.E. CROSS, W.T. TACK, L.W. LOECHEL, AND L.S. KRAMER, ALUMINUM WELDABILITY AND HOT TEARING THEORY, WELDABILITY OF MATERIALS, ASM INTERNATIONAL, 1990, P 275 9. R.A. CHIHOSKI, THE CHARACTER OF STRESS FIELDS AROUND A WELD ARC MOVING ON ALUMINUM SHEET, WELD J., VOL 51 (NO. 1), 1972, P 9-S 10. C.L. CROSS, "WELDABILITY OF ALUMINUM-LITHIUM ALLOYS: AN INVESTIGATION OF HOT TEARING MECHANISMS," PH.D. THESIS, COLORADO SCHOOL OF MINES, GOLDEN, CO, 1986, P 144 11. A.R.E. SINGER AND P.H. JENNINGS, HOT SHORTNESS OF THE ALUMINUM SILICON ALLOYS OF COMMERCIAL PURITY, J. INST. MET., VOL 73, 1947, P 197 12. W.I. PUMPHREY AND J.V. LYONS, CRACKING DURING THE CASTING AND WELDING OF THE MORE COMMON BINARY ALUMINUM ALLOYS, J. INST. MET., VOL 74, 1948, P 439 13. J.D. DOWD, WELD CRACKING OF ALUMINUM ALLOYS, WELD. J., VOL 31 (NO. 10), 1952, P 448S 14. P.H. JENNINGS, A.R.E. SINGER, AND W.I. PUMPHREY, HOT-SHORTNESS OF SOME HIGHPURITY ALLOYS IN THE SYSTEMS ALUMINUM-COPPER-SILICON AND ALUMINUMMAGNESIUM-SILICON, J. INST. MET., 1948, P 227 15. J.H. DUDAS AND F.R. COLLINS, PREVENTING CRACKS IN HIGH-STRENGTH ALUMINUM ALLOYS, WELD. J., VOL 45 (NO. 6), 1966, P 241-S 16. W.I. PUMPHREY AND D.C. MOORE, CRACKING DURING AND AFTER SOLIDIFICATION IN SOME ALUMINUM-COPPER-MAGNESIUM ALLOYS OF HIGH PURITY, J. INST. MET., VOL 74, 1948, P 425 17. J.K. DAWSON, WELD CRACK SENSITIVITY OF ALUMINUM ALLOYS, MET. PROG., VOL 76 (NO. 1), 1959, P 116

18. R.P. MEISTER AND D.C. MARTIN, "WELDING OF ALUMINUM AND ALUMINUM ALLOYS," DEFENSE METALS INFORMATION CENTER, BATTELLE MEMORIAL INSTITUTE, 1967, P 51 19. S. KOU, WELDING METALLURGY, JOHN WILEY & SONS, 1987, P 239 20. R.P. MARTUKANITZ, ALCOA LABORATORIES, UNPUBLISHED RESEARCH, 1985 21. N.F. GITTOS AND M.H. SCOTT, HEAT AFFECTED ZONE CRACKING OF AL-MG-SI ALLOYS, WELD. J., VOL 60 (NO. 6), 1981, P 95-S 22. H.W. KERR AND M. KATOH, INVESTIGATION OF HEAT-AFFECTED ZONE CRACKING OF GMA WELDS OF AL-MG-SI ALLOYS USING THE VARESTRAINT TEST, WELD. J., VOL 66 (NO. 9), 1987, P 251-S 23. D.G. HOWDEN AND D.R. MILNER, BRIT. WELD. J., VOL 10, 1963, P 304 24. Z.P. SAPERSTEIN, G.R. PRESCOTT, AND E.W. MONROE, POROSITY IN ALUMINUM WELDS, WELD. J., VOL 43 (NO. 10), 1964, P 443-S 25. C. RANSLEY AND J. NEUFELD, J. INST. MET., VOL 74, 1948, P 617 26. M. UDA AND S. OHNO, TRANS. NATL. RES. INST. MET. (JPN), VOL 16, 1974, P 2 27. R.P. MARTUKANITZ, ALCOA LABORATORIES, UNPUBLISHED RESEARCH, 1986 28. R.P. MARTUKANITZ AND P.R. MICHNUK, SOURCES OF POROSITY IN GAS METAL ARC WELDING OF ALUMINUM, ALUMINIUM, VOL 58 (NO. 5), 1982, P 276 29. P.B. DICKERSON, ALCOA LABORATORIES, PRIVATE COMMUNICATION, 1987 30. A.N. O'SHANSKII AND V.V. D'YACHENCHO, EVALUATION OF THE SUSCEPTIBILITY OF ALLOYS TO THE FORMATION OF PORES IN WELDING, WELD. PROD., VOL 24 (NO. 7), 1977, P 52 31. S.D. DUMOLT, "METALLURGICAL TRANSACTIONS IN THE HEAT AFFECTED ZONE OF ALUMINUM ALLOYS BY TRANSMISSION ELECTRON MICROSCOPY," PH.D. THESIS, CARNEGIE-MELLON UNIVERSITY, PITTSBURGH, PA, 1983 32. T. ENGO AND T. KURODA, MICROSTRUCTURE IN WELD HEAT-AFFECTED ZONE OF AL-MGSI ALLOYS, TRANS. JPN. WELD. RES. INST., VOL 11 (NO. 1), 1982, P 61 33. R.P. MARTUKANITZ, "MODELING THE HEAT AFFECTED ZONES OF ALUMINUM ARC WELDS, ADVANCES IN WELDING SCIENCE AND TECHNOLOGY, S.A. DAVID, ED., ASM INTERNATIONAL, 1986, P 193 34. S. KOU, WELDING METALLURGY AND WELDABILITY OF HIGH STRENGTH ALUMINUM ALLOYS, WELD. RES. COUNC. BULL., NO. 320, WELDING RESEARCH COUNCIL, 1986 35. R.P. MARTUKANITZ, ALCOA LABORATORIES, UNPUBLISHED RESEARCH, 1985 36. S.D. DUMOLT, PH.D. THESIS, CARNEGIE-MELLON UNIVERSITY, PITTSBURGH, PA, 1983 37. W.L. BURCH, THE EFFECT OF WELDING SPEED ON STRENGTH OF 6061-T4 ALUMINUM JOINTS, WELD. J., VOL 37 (NO. 7), 1958, P 361-S 38. R.P. MARTUKANITZ, "MODELING THE HEAT AFFECTED ZONES OF ALUMINUM ARC WELDS, ADVANCES IN WELDING SCIENCE AND TECHNOLOGY, S.A. DAVID, ED., ASM INTERNATIONAL, 1986, P 195 39. R.J. BRUNGRABER AND F.G. NELSON, EFFECT OF WELDING VARIABLES ON ALUMINUM ALLOY WELDMENTS, WELD. J., VOL 52 (NO. 3), 1973, P 97-S 40. G.E. METZGER, SOME MECHANICAL PROPERTIES OF WELDS IN 6061 ALUMINUM ALLOY SHEET, WELD. J., VOL 56 (NO. 10), 1967, P 457-S Selection and Weldability of Heat-Treatable Aluminum Alloys Richard P. Martukanitz, Applied Research Laboratory, Pennsylvania State University

Filler Alloy Selection and Properties of Welds

Although the selected base alloy and welding process greatly influence the characteristics of aluminum joints, the filler alloy plays a significant role in establishing a number of important weld properties for hem-treatable aluminum alloys. Of foremost concern is the weldability of a particular base alloy and filler alloy combination, which is defined as its susceptibility to weld cracking. Shear strength of a fillet weld, response to postweld heat treatments, and color match after anodizing are other attributes that are highly governed by the filler alloy used to produce the weld in heat-treatable aluminum alloys. Properties such as joint strength, ductility, toughness, and corrosion resistance across the weld are also dictated by the base alloy and filler alloy combination selected. The filler alloy selection chart in the article "Welding of Aluminum Alloys" in this Volume (see Table 14 in that article) provides a comprehensive guide for selecting the proper filler metal. Recently a knowledge-based system using computer software has also been constructed to allow greater interaction in selecting aluminum base metal and filler metal alloys for welding (Ref 41). Filler Alloy Selection and Weld Crack Sensitivity Weld crack sensitivity is most easily controlled by influencing the composition of the weld. This is achieved by selecting the proper filler alloy and ensuring adequate dilution of the filler alloy to the weld. Generally, joint designs that increase filler metal dilution (for example, square groove with gap as opposed to no gap) are beneficial in reducing the tendency for cracking. The principal filler alloys for welding aluminum are the 1xxx (1100), 2xxx (2319), 4xxx (4043, 4047, 4145, and 4643), and 5xxx (5154, 5183, 5356, 5554, 5556, and 5654) series. Nominal compositions of aluminum filler alloys are shown in Table 4 (Ref 42). Most of these alloys are high in solute to reduce crack sensitivity. Because of their narrow solidification temperature range, the 4xxx series filler alloys provide excellent insensitivity to weld cracking but are not applicable for welding all aluminum-base alloys. Because of the formation of large amounts of brittle magnesium-silicide (Mg2Si), the 4xxx series filler alloys are not applicable for welding the 7xxx series alloys containing appreciable amounts of magnesium. Table 5 shows the filler alloys recommended for welding various base alloys to minimize sensitivity to weld cracking.

TABLE 4 NOMINAL COMPOSITION OF ALUMINUM FILLER METAL ALLOYS

ALLOY COMPOSITION, WT% Si Cu Mn Mg Cr 1100 . . . 0.12 . . . . . . . . . 1188 ... ... ... ... ... 2319 . . . 6.3 0.30 . . . . . . 4043 5.2 . . . . . . . . . . . . 4047 12.0 . . . . . . . . . . . . 4145 10.0 4.0 . . . . . . . . . 4643 4.1 . . . . . . 0.2 . . . 5183 . . . . . . 0.75 4.7 0.15 5356 . . . . . . 0.12 5.0 0.12 5554 . . . . . . 0.75 2.7 0.12 5556 . . . . . . 0.75 5.1 0.12 5654 . . . . . . . . . 3.5 0.25 C355.0 5.0 1.25 . . . 0.5 . . . A356.0 7.0 . . . . . . 0.4 . . . 357.0 7.0 . . . . . . 0.5 . . . A357.0 7.0 . . . . . . 0.6 . . . Source: Ref 42

Ti ... ... 0.15 ... ... ... ... ... 0.13 0.12 0.12 0.10 ... ... ... 0.12

OTHER 99.00 AL MIN 99.88 AL MIN 0.17 ZR, 0.10 V ... ... ... ... ... ... ... ... ... ... ... ... ...

TABLE 5 FILLER ALLOYS RECOMMENDED FOR WELDING SELECTED ALUMINUM-BASE ALLOYS TO MINIMIZE SENSITIVITY TO WELD CRACKING BASE METAL

1060, 1070, 1080, 1350 1100, 3003, ALCLA D 3003 2014, 2036 2219, 2519 3004, ALCLA D 3004 5005, 5050 5052, 5652(B) 5083

319.0, 333.0, 354.0, 355.0, C355.0, 380.0

356.0, A356.0, A357.0, 359.0, 413.0, A444,0, 443.0

511.0, 512.0, 513.0, 514.0

7005(A), 7039, 710.0, 711.0, 712.0

6070

4145(C)(D)

4043(D)(C)

4043(D)

4043(D)

4043(D)

6061, 6463, 6101, 6201, 6151, 6351, 6951

5456

4043(D)

5356(

(F)

4145(C)(D)

4043(D)(E)

4043(D)

4043(D)

4043(D)

4043(D)

4145

...

...

4145

4145

5086

5083

5052, 5652(B)

5005 5050

3004 ALCL AD 3004

2219, 2519

2014 , 2036

1100, 1060, 3004, ALCL AD 3003

1060 1070 , 1080 , 1350

4043(D)

4043(D)

5356(

5356(

4043(D)

1100(C)

4043

4145

4145

1100(C)

1188(

(F)

C)

C)

G)

5356(

4043(D)

4043(D)

5356(

5356(

4043(D)(F

C)

(F)

(F)

C)

C)

)

...

...

...

...

...

...

4043(F)

4043(F)

4145

4145

1100(C)

...

...

...

4145(H)

4145(

...

...

...

...

...

H)

4145(C)(D)

4145(C)(D)

4043(D)

4043(D)

(H)

4043(D)

4043(D)

5654(I)

5356(F)

4043(D)

4043(D)

(E)

(E)

4043(F)

4043(I)

4043

4043(D)

4043(D)

4043(D)

4043(D)

5654(I)

5356(F)

4043(F)

4043(I)

5356(

5654(I)

4043(D)(I) 5356(C)(D)

5654(I) 5356(F)

5356(F) 5183(F)

5356(C)(D)

5086

5356(

5356(C)(

5356(C)(

5356(

I)

I)

I)

5356(F)

5356(F)

5183(

5356(F)

5356(F)

5356(F)

5356(F)

043(D)(I) 4043(D)

4043(D)(I) 4043(D)(I)

5456

5654(I) 5356(F)

5356(I) 5356(I) 5556(F)

4043

4043

2319(C)(D) ...

...

...

...

5654(I)

4043(D)(F

4043(F)

...

...

...

...

...

)

(J)

5356(

5654(B)(C

...

...

...

...

...

...

F)

)(I)

5356(

5183(

...

...

...

...

...

...

...

F)

F)

5356(

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

5654(I)

5654(I) 5356(F)

5654(I) 5356(F)

5356(

5356(I)

5356(I)

F)

5654(I)

4043(D)

4043(F)

F)

(F)

4043

4043(F)

5654(I)

F)

4043(D)

4043

(E)

F)

(F)

5154, 5254(B) 5454

5154, 5254(B )

C)

(F)

4145(H)

5454

5356(

5356(

4043(D)(F

F)

F)

)

5356(

5356(

F)

F)

5356( F)

F)

5356(C)(

5356(C)(

5356(

I)

I)

I)

5654(I)

5354(B)

5356(C)(

5356(C)(

5356(

5554(C)

I)

I)

I)

(F)

5356(F)

5356(F)

5556(

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

(I)

F)

6061,60 63, 6351, 6101, 6201, 6151, 6951 6070

4145(C)(D)

4145(C)(D)

5356(C)(D)

5356(C)

5356(C)(D)

4043(D)

4043(D)

(F)

(I)

(I)

(I)

(I)

4043(D)(F)

5356(C)

5356(C)(D)

4043(D)

...

(F) (A)

7005 , 7039, 710.0, 711.0, 712.0 511.0, 512.0, 513.0, 514.0 356.0, A356.0, A357.0, 359.0, 413.0, A444.0, 443.0 319.0, 333.0, 354.0, 355.0, C355.0 380.0

(D)

4043

(D)(I)

(F) (I)

4043

5356

5356

4043(D)(I)

5654(I)(

(F) (F)

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

J)

4145(C)(D)

4043(D)(J)

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

4145(C)(D)

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

...

(J)

Notes: (1) Service conditions such as immersion in fresh or salt water, exposure to specific chemicals, or a sustained high temperature (higher than 65 °C, or 150 °F) may limit choice of filler metals Filler alloys 5356, 5183, 5556, and 5654 are not recommended for sustained elevated-temperature service. (2) Recommendations in this table apply to gas-shielded arc welding processes. For gas welding, only 1100, 1188, and 4043 filler metals are ordinarily used. (3) All filler metals are listed in AWS specification A5.10. (4) Where no filler metal is listed, the base metal combination is not recommended for welding. Source: Ref 43

(A) REFERS TO 7005 EXTRUSIONS ONLY (X-PREFIX STILL APPLIES TO SHEET AND PLATE). (B) BASE METAL ALLOYS 5652 AND 5654 ARE USED FOR HYDROGEN PEROXIDE SERVICE. 5654 FILLER METAL IS USED FOR WELDING BETH ALLOYS FOR LOW-TEMPERATURE SERVICE (65 °C, OR 150 °F, OR BELOW). (C) 4043 MAY BE USED FOR SOME APPLICATIONS. (D) 4047 MAY BE USED FOR SOME APPLICATIONS. (E) 4145 MAY BE USED FOR SOME APPLICATIONS. (F) 5183, 5356, OR 5556 MAY BE USED. (G) 1100 MAY BE USED FOR SOME APPLICATIONS. (H) 2319 MAY BE USED FOR SOME APPLICATIONS. (I) 5183, 5356, 5554, 5556. OR 5654 MAY BE USED. IN SOME CASES THEY PROVIDE IMPROVED COLOR MATCH AFTER ANODIZING TREATMENT, HIGHEST WELD QUALITY, AND HIGHER WELD STRENGTH 5554 IS SUITABLE FOR ELEVATED TEMPERATURE SERVICE.

(J) FILLER METAL WITH THE SAME ANALYSIS AS THE BASE METAL IS SOMETIMES USED.

Weld Strength and Ductility The strength and ductility of welds in aluminum alloys are determined by a number of parameters. These may include base alloy and temper, filler alloy, welding process and parameters, use of postweld heat treatments, joint type, principal loading condition, and the presence of weld defects. As described earlier, metallurgical transformations in the HAZ of heat-treatable aluminum alloys result in a degradation of strength in this region. The degree of degradation may be related to heat input used to produce the weld. Processes that provide higher travel speeds minimize the degradation in the HAZ and typically result in higher weld strength. Postweld heat treatments may be used to recover the loss of strength in the HAZ. Postweld solution heat treating and aging provides the greatest improvement in joint strength; however, this practice requires the use of a cold-water quench that may impart unacceptable part distortion. Postweld aging, which is a lower-temperature treatment, results in a moderate improvement in joint strength and does not require water quenching. It is also appropriate to comment on the effects of loading that distinguish the properties of welded joints. The transverse tensile strength of a butt weld provides a relative indication of joint strength in service, although butt welds may also undergo multiaxial loading that hinders direct comparison to tabular data representing uniaxial behavior. In the case of fillet welds, the transverse or longitudinal shear strength (depending on the principal loading condition) obtained from shear test specimens may be used to ascertain joint performance. The strength and ductility of groove welds in the aswelded condition, the postweld aged only condition, and the solution heat treated and aged (artificially and naturally aged) condition are described in the article "Welding of Aluminum Alloys" in this Volume. Transverse and longitudinal shear strengths of fillet welds are also listed in the aforementioned article. Additional information on properties of aluminum weldments can be found in Ref 44 and 45. Figure 12 displays strength profiles across gas-tungsten arc welds of 2219-T87 and 6061-T6 alloys. The 2219-T87 alloy was welded with 2319 filler alloy and the 6061-T6 was welded with 4043 alloy. The strength profile of the workhardenable alloy 5456-H116 welded with filler alloy 5356 is also shown in the figure for comparison. The strength values in Fig. 12 are approximations obtained by converting hardness measurements using the empirically derived relationship for aluminum alloys (Ref 46):  HRE  UTS = 10, 66 − 19, 42 ln  1 −  109  

(EQ 2)

where UTS is the ultimate tensile strength in ksi, and HRE is the Rockwell hardness using the E-scale.

FIG. 12 STRENGTH PROFILES ACROSS ALTERNATING CURRENT GAS-TUNGSTEN ARC WELDS ON 3.2 MM (0.125 IN.) THICK 2219-T87 USING 2319 FILLER ALLOY, 6061-T6 USING 4043 FILLER ALLOY, AND 5456-H116 USING 5356 FILLER ALLOY. SOURCE: REF 35

It is evident from Fig. 12 that the degradation of strength associated with the HAZ usually dictates the strength of the joint for aluminum alloys. This is not true when weldments have been postweld heat treated, because the fusion zone may not respond to these treatments to the same degree as the HAZ. The 2xxx series alloys exhibit the lowest ductility of the precipitation-hardenable alloys. This is due to the formation of brittle, copper-bearing eutectic phases at grain boundaries near the fusion zone interface. The 6xxx and 7xxx series alloys display good weld ductility. Postweld solution heat treating and aging of the heat-treat-able alloys or postweld aging only, usually decreases weld ductility. Corrosion Resistance of Welds Many of the heat-treatable aluminum alloys can be welded without reducing the resistance of the assembly to corrosion. However, because welding generates residual tensile stresses and modifies the parent metal structure, some alloys may have their resistance to corrosion lowered by welding. Corrosion of aluminum weldments is usually localized or preferential in nature and depends on the base alloy, filler alloy, and HAZ structure. The corrosion potential variations determined by electrode potential measurements across 2219 and 7039 weldments are shown in Fig. 13. The corrosion potential variations across the weld of the work-hardenable alloy 5456 are also shown in Fig. 13. The corrosion potentials of various base and filler alloys are also described in the article "Welding of Aluminum Alloys" in this Volume.

FIG. 13 PLOT OF CORROSION POTENTIAL AND HARDNESS LEVELS VERSUS DISTANCE FROM WELD CENTERLINE FOR SELECTED ALUMINUM ALLOYS. (A) ALLOY 5456-H321 (5556 FILLER METAL; DCRP-MIG WELD PRODUCED IN THREE PASSES). (B) 2219-T87 (2319 FILLER METAL; DCSP-TIG WELD PRODUCED IN TWO PASSES). (C) 7039-T651 (5183 FILLER METAL; DCRP-TIG WELD PRODUCED IN TWO PASSES). DCRP, DIRECT CURRENT REVERSE POLARITY; DCSP, DIRECT CURRENT STRAIGHT POLARITY. COURTESY OF ALCOA LABORATORIES

The 6xxx series alloys are highly resistant to localized corrosion after welding. The weldable alloy 2219 of the 2xxx series also exhibits good corrosion resistance after welding. The corrosion potential for 2219-T87 weldments produced using 2319 filler metal is shown in Fig. 13 and depicts a lower corrosion potential in the weld area as compared to the unaffected base metal. The weldable 7xxx series alloys produce greater sensitivity to corrosion after welding (Ref 47). In this case, the structural transformation of the HAZ due to the heat of welding makes the HAZ anodic to either the 5xxx

weld metal or the unaffected base metal. This mechanism is also shown in Fig. 13 for 7089-T651 welded with 5183 filler metal. The highly anodic, small area of the HAZ will pit and corrode in the presence of an electrolyte to protect both the weld metal and the much larger unaffected base metal. Postweld solution heat treating and aging of the 7xxx series alloys will improve their corrosion resistance; however, aluminum-magnesium filler alloys containing more than 3.5% Mg should not be used when postweld heat treating because the fusion zone can be sensitized to stress-corrosion cracking (Ref 48). Wrought alloys usually have greater resistance to stress-corrosion cracking in the longitudinal orientation (direction of working) than in the transverse orientation or in the short-transverse orientation (through the thickness). Because of this, welding of the 7xxx series alloys near a base-metal edge may result in a tensile stress in the short-transverse direction sufficient to cause stress-corrosion cracking in the exposed edge. "Buttering" the edge with weld metal provides compressive stress at the edge and overcomes the stress-corrosion cracking problem. Finally, paint protection or metallizing with selected aluminum alloys can be used to reduce the general corrosion sensitivity of heat-treated aluminum weldments. Other Joint Characteristics Other joint characteristics may be of importance, depending upon the service requirements of the welded assembly. These may include characteristics that can affect the structural performance of the assembly, such as strength and toughness at the service temperature, as well as aesthetic qualities, such as joint color match after anodizing. Aluminum alloys are compatible with all known cryogenics; base-metal strength and weld strength generally increase at cryogenic temperatures (Ref 49 and 50). The 2xxx series heat-treatable alloys, as well as the work-hardenable 5xxx series alloys, display a good combination of strength and ductility at low temperatures. Alloy 2219 is used extensively for cryogenic applications requiring high strength and good weldability. Toughness of welds in aluminum alloys is defined as the ability of the joint to absorb energy by deformation rather than fracture catastrophically. Fracture toughness in aluminum alloy welds is typically not determined by Charpy impact tests but may be reasonably obtained by tear or tension tests involving notched specimens (Ref 51, 52, 53). Based on results of these tests, welds of 2219 and 7039 possess good notch toughness, although it must also be mentioned that welds representing the work-hardenable alloys of the 5xxx series also exhibit good toughness. Although the weld and base-metal strengths of the heat-treatable aluminum alloys, particularly the 2xxx series, fare well at elevated temperatures up to 150 °C (300 °F), caution must be observed when selecting filler alloys for welded assemblies functioning at these temperatures. Elevated-temperature service for aluminum welds is usually defined as service at sustained temperatures above 65 °C (150 °F). Sustained elevated-temperature service of welds containing high amounts of magnesium may result in increased sensitivity to stress-corrosion cracking due to the formation of Mg2Al3 at grain boundaries. Therefore, most welded applications involving sustained elevated-temperature service dictate the use of filler alloys that do not contain magnesium. An exception is the use of filler alloy 5554 (containing approximately 2.7% Mg) when welds of 6xxx series alloys at sustained elevated temperatures also require good color match after anodizing. Typically, 2319 filler alloy is used in welds on 2xxx series base alloys that require good color match after anodizing. When good color match after anodizing is desired for welds in the 6xxx or 7xxx series base alloys, filler alloys of the 5xxx series containing magnesium are utilized. High levels of silicon in the weld when using the 4xxx series filler alloys result in darkening of the fusion zone in comparison to the base metal after anodizing.


35. R.P. MARTUKANITZ, ALCOA LABORATORIES, UNPUBLISHED RESEARCH, 1985 41. R.J. NAPOLITANO, H. KULLUK, M.L. NAGURKA, R.P. MARTUKANITZ, AND P.B. DICKERSON, DEVELOPMENT OF KNOWLEDGE-BASED SYSTEMS FOR ALUMINUM WELDING, WELD. J., VOL 71 (NO. 12), 1992, P 43 42. P.B. DICKERSON AND B. IRVING, WELDING ALUMINUM: IT'S NOT AS DIFFICULT AS IT SOUNDS, WELD. J., VOL 71 (NO. 4), 1992, P 47 43. WELDING ALUMINUM: THEORY AND PRACTICE, H.L. SAUNDERS, ED., THE ALUMINUM ASSOCIATION, 1991, P 3.8 44. WELDING ALUMINUM, AMERICAN WELDING SOCIETY, 1972, P 69.94-69.106 45. F.G. NELSON AND R.L. ROLF, SHEAR STRENGTH OF AL ALLOY FILLET WELDS, WELD. J.,

VOL 45 (NO. 2), 1966, P 82-S 46. B.S. SHABEL AND R.F. YOUNG, PROC. CONF. NONDESTRUCTIVE CHARACTERIZATION OF MATERIALS II, J.F. BUSSIERE, J.P. MONCHALIN, C.O. RUUD, AND R.E. GREEN, JR., ED., PLENUM, 1987, P 335 47. R.P. MEISTER AND D.C. MARTIN, "WELDING OF ALUMINUM AND ALUMINUM ALLOYS," DMIC REPORT 236, DEFENSE METALS INFORMATION CENTER, 1967, P 65 48. "ALCOA ALUMINUM ALLOY 7005 GREEN LETTER," ALUMINUM COMPANY OF AMERICA, 1974 49. R.P. MEISTER AND D.C. MARTIN, "WELDING OF ALUMINUM AND ALUMINUM ALLOYS," DMIC REPORT 236, DEFENSE METALS INFORMATION CENTER, 1967, P 55 50. J.F. CAMPBELL, ALUMINUM ALLOYS FOR CRYOGENIC SERVICE, MATER. RES. STAND., VOL 4 (NO. 10), 1964, P 540 51. F.G. NELSON AND J.G. KAUFMAN, FRACTURE CHARACTERISTICS OF WELDS IN ALUMINUM ALLOYS, WELD. J., VOL 45 (NO. 7), 1966, P 321-S 52. ALUMINUM FOR CRYOGENIC ENVIRONMENT, PART I: GENERAL CONSIDERATIONS, CRYOG. ENG. NEWS, BUSINESS COMMUNICATIONS PUBLISHERS, MARCH 1966 53. ALUMINUM FOR CRYOGENIC ENVIRONMENT, PART II: ALLOY GROUPS, CRYOG. ENG. NEWS, BUSINESS COMMUNICATIONS PUBLISHERS, APRIL 1966 Selection and Weldability of Heat-Treatable Aluminum Alloys Richard P. Martukanitz, Applied Research Laboratory, Pennsylvania State University

References

1. P.B. DICKERSON AND B. IRVING, WELDING ALUMINUM: IT'S NOT AS DIFFICULT AS IT SOUNDS, WELD. J., VOL 71 (NO. 4), 1992, P 45 2. ALUMINUM: PROPERTIES AND PHYSICAL METALLURGY, J.E. HATCH, ED., AMERICAN SOCIETY FOR METALS, 1984, P 222 3. ALUMINUM STANDARDS AND DATA, THE ALUMINUM ASSOCIATION, INC., 1988, P 15 4. ALUMINUM: PROPERTIES AND PHYSICAL METALLURGY, J.E. HATCH, ED., AMERICAN SOCIETY FOR METALS, 1984, P 320 5. R.S. JAMES, ALUMINUM-LITHIUM ALLOYS, VOL 2, ASM HANDBOOK, ASM INTERNATIONAL, 1990, P 178-199 6. ALUMINUM STANDARDS AND DATA, THE ALUMINUM ASSOCIATION, INC., 1988, P 11 7. J.M. SILCOCK, T.J. HEAL, AND H.K. HARDY, STRUCTURAL AGEING CHARACTERISTICS OF BINARY ALUMINUM-COPPER ALLOYS, J. INST. MET., VOL 82, 1953-1954, P 239 8. C.E. CROSS, W.T. TACK, L.W. LOECHEL, AND L.S. KRAMER, ALUMINUM WELDABILITY AND HOT TEARING THEORY, WELDABILITY OF MATERIALS, ASM INTERNATIONAL, 1990, P 275 9. R.A. CHIHOSKI, THE CHARACTER OF STRESS FIELDS AROUND A WELD ARC MOVING ON ALUMINUM SHEET, WELD J., VOL 51 (NO. 1), 1972, P 9-S 10. C.L. CROSS, "WELDABILITY OF ALUMINUM-LITHIUM ALLOYS: AN INVESTIGATION OF HOT TEARING MECHANISMS," PH.D. THESIS, COLORADO SCHOOL OF MINES, GOLDEN, CO, 1986, P 144 11. A.R.E. SINGER AND P.H. JENNINGS, HOT SHORTNESS OF THE ALUMINUM SILICON ALLOYS OF COMMERCIAL PURITY, J. INST. MET., VOL 73, 1947, P 197 12. W.I. PUMPHREY AND J.V. LYONS, CRACKING DURING THE CASTING AND WELDING OF THE MORE COMMON BINARY ALUMINUM ALLOYS, J. INST. MET., VOL 74, 1948, P 439

13. J.D. DOWD, WELD CRACKING OF ALUMINUM ALLOYS, WELD. J., VOL 31 (NO. 10), 1952, P 448S 14. P.H. JENNINGS, A.R.E. SINGER, AND W.I. PUMPHREY, HOT-SHORTNESS OF SOME HIGHPURITY ALLOYS IN THE SYSTEMS ALUMINUM-COPPER-SILICON AND ALUMINUMMAGNESIUM-SILICON, J. INST. MET., 1948, P 227 15. J.H. DUDAS AND F.R. COLLINS, PREVENTING CRACKS IN HIGH-STRENGTH ALUMINUM ALLOYS, WELD. J., VOL 45 (NO. 6), 1966, P 241-S 16. W.I. PUMPHREY AND D.C. MOORE, CRACKING DURING AND AFTER SOLIDIFICATION IN SOME ALUMINUM-COPPER-MAGNESIUM ALLOYS OF HIGH PURITY, J. INST. MET., VOL 74, 1948, P 425 17. J.K. DAWSON, WELD CRACK SENSITIVITY OF ALUMINUM ALLOYS, MET. PROG., VOL 76 (NO. 1), 1959, P 116 18. R.P. MEISTER AND D.C. MARTIN, "WELDING OF ALUMINUM AND ALUMINUM ALLOYS," DEFENSE METALS INFORMATION CENTER, BATTELLE MEMORIAL INSTITUTE, 1967, P 51 19. S. KOU, WELDING METALLURGY, JOHN WILEY & SONS, 1987, P 239 20. R.P. MARTUKANITZ, ALCOA LABORATORIES, UNPUBLISHED RESEARCH, 1985 21. N.F. GITTOS AND M.H. SCOTT, HEAT AFFECTED ZONE CRACKING OF AL-MG-SI ALLOYS, WELD. J., VOL 60 (NO. 6), 1981, P 95-S 22. H.W. KERR AND M. KATOH, INVESTIGATION OF HEAT-AFFECTED ZONE CRACKING OF GMA WELDS OF AL-MG-SI ALLOYS USING THE VARESTRAINT TEST, WELD. J., VOL 66 (NO. 9), 1987, P 251-S 23. D.G. HOWDEN AND D.R. MILNER, BRIT. WELD. J., VOL 10, 1963, P 304 24. Z.P. SAPERSTEIN, G.R. PRESCOTT, AND E.W. MONROE, POROSITY IN ALUMINUM WELDS, WELD. J., VOL 43 (NO. 10), 1964, P 443-S 25. C. RANSLEY AND J. NEUFELD, J. INST. MET., VOL 74, 1948, P 617 26. M. UDA AND S. OHNO, TRANS. NATL. RES. INST. MET. (JPN), VOL 16, 1974, P 2 27. R.P. MARTUKANITZ, ALCOA LABORATORIES, UNPUBLISHED RESEARCH, 1986 28. R.P. MARTUKANITZ AND P.R. MICHNUK, SOURCES OF POROSITY IN GAS METAL ARC WELDING OF ALUMINUM, ALUMINIUM, VOL 58 (NO. 5), 1982, P 276 29. P.B. DICKERSON, ALCOA LABORATORIES, PRIVATE COMMUNICATION, 1987 30. A.N. O'SHANSKII AND V.V. D'YACHENCHO, EVALUATION OF THE SUSCEPTIBILITY OF ALLOYS TO THE FORMATION OF PORES IN WELDING, WELD. PROD., VOL 24 (NO. 7), 1977, P 52 31. S.D. DUMOLT, "METALLURGICAL TRANSACTIONS IN THE HEAT AFFECTED ZONE OF ALUMINUM ALLOYS BY TRANSMISSION ELECTRON MICROSCOPY," PH.D. THESIS, CARNEGIE-MELLON UNIVERSITY, PITTSBURGH, PA, 1983 32. T. ENGO AND T. KURODA, MICROSTRUCTURE IN WELD HEAT-AFFECTED ZONE OF AL-MGSI ALLOYS, TRANS. JPN. WELD. RES. INST., VOL 11 (NO. 1), 1982, P 61 33. R.P. MARTUKANITZ, "MODELING THE HEAT AFFECTED ZONES OF ALUMINUM ARC WELDS, ADVANCES IN WELDING SCIENCE AND TECHNOLOGY, S.A. DAVID, ED., ASM INTERNATIONAL, 1986, P 193 34. S. KOU, WELDING METALLURGY AND WELDABILITY OF HIGH STRENGTH ALUMINUM ALLOYS, WELD. RES. COUNC. BULL., NO. 320, WELDING RESEARCH COUNCIL, 1986 35. R.P. MARTUKANITZ, ALCOA LABORATORIES, UNPUBLISHED RESEARCH, 1985 36. S.D. DUMOLT, PH.D. THESIS, CARNEGIE-MELLON UNIVERSITY, PITTSBURGH, PA, 1983 37. W.L. BURCH, THE EFFECT OF WELDING SPEED ON STRENGTH OF 6061-T4 ALUMINUM JOINTS, WELD. J., VOL 37 (NO. 7), 1958, P 361-S 38. R.P. MARTUKANITZ, "MODELING THE HEAT AFFECTED ZONES OF ALUMINUM ARC WELDS, ADVANCES IN WELDING SCIENCE AND TECHNOLOGY, S.A. DAVID, ED., ASM

INTERNATIONAL, 1986, P 195 39. R.J. BRUNGRABER AND F.G. NELSON, EFFECT OF WELDING VARIABLES ON ALUMINUM ALLOY WELDMENTS, WELD. J., VOL 52 (NO. 3), 1973, P 97-S 40. G.E. METZGER, SOME MECHANICAL PROPERTIES OF WELDS IN 6061 ALUMINUM ALLOY SHEET, WELD. J., VOL 56 (NO. 10), 1967, P 457-S 41. R.J. NAPOLITANO, H. KULLUK, M.L. NAGURKA, R.P. MARTUKANITZ, AND P.B. DICKERSON, DEVELOPMENT OF KNOWLEDGE-BASED SYSTEMS FOR ALUMINUM WELDING, WELD. J., VOL 71 (NO. 12), 1992, P 43 42. P.B. DICKERSON AND B. IRVING, WELDING ALUMINUM: IT'S NOT AS DIFFICULT AS IT SOUNDS, WELD. J., VOL 71 (NO. 4), 1992, P 47 43. WELDING ALUMINUM: THEORY AND PRACTICE, H.L. SAUNDERS, ED., THE ALUMINUM ASSOCIATION, 1991, P 3.8 44. WELDING ALUMINUM, AMERICAN WELDING SOCIETY, 1972, P 69.94-69.106 45. F.G. NELSON AND R.L. ROLF, SHEAR STRENGTH OF AL ALLOY FILLET WELDS, WELD. J., VOL 45 (NO. 2), 1966, P 82-S 46. B.S. SHABEL AND R.F. YOUNG, PROC. CONF. NONDESTRUCTIVE CHARACTERIZATION OF MATERIALS II, J.F. BUSSIERE, J.P. MONCHALIN, C.O. RUUD, AND R.E. GREEN, JR., ED., PLENUM, 1987, P 335 47. R.P. MEISTER AND D.C. MARTIN, "WELDING OF ALUMINUM AND ALUMINUM ALLOYS," DMIC REPORT 236, DEFENSE METALS INFORMATION CENTER, 1967, P 65 48. "ALCOA ALUMINUM ALLOY 7005 GREEN LETTER," ALUMINUM COMPANY OF AMERICA, 1974 49. R.P. MEISTER AND D.C. MARTIN, "WELDING OF ALUMINUM AND ALUMINUM ALLOYS," DMIC REPORT 236, DEFENSE METALS INFORMATION CENTER, 1967, P 55 50. J.F. CAMPBELL, ALUMINUM ALLOYS FOR CRYOGENIC SERVICE, MATER. RES. STAND., VOL 4 (NO. 10), 1964, P 540 51. F.G. NELSON AND J.G. KAUFMAN, FRACTURE CHARACTERISTICS OF WELDS IN ALUMINUM ALLOYS, WELD. J., VOL 45 (NO. 7), 1966, P 321-S 52. ALUMINUM FOR CRYOGENIC ENVIRONMENT, PART I: GENERAL CONSIDERATIONS, CRYOG. ENG. NEWS, BUSINESS COMMUNICATIONS PUBLISHERS, MARCH 1966 53. ALUMINUM FOR CRYOGENIC ENVIRONMENT, PART II: ALLOY GROUPS, CRYOG. ENG. NEWS, BUSINESS COMMUNICATIONS PUBLISHERS, APRIL 1966

Selection and Weldability of Non-Heat-Treatable Aluminum Alloys C.E. Cross, Martin Marietta Astronautics Group; M.L. Kohn, FMC Corporation

Introduction NON-HEAT-TREATABLE ALUMINUM ALLOYS constitute a group of alloys that rely solely upon cold work and solid solution strengthening for their strength properties. They differ from heat-treatable alloys in that they are incapable of forming second-phase precipitates for improved strength (see the article "Selection and Weldability of Heat-Treatable Aluminum Alloys" in this Volume). As such, the non-heat-treatable alloys cannot achieve the high strengths characteristic of precipitation-hardened alloys. For example, the yield strength of precipitation-hardened alloy 7075-T6 (505 MPa, or 73 ksi) is significantly greater than that of cold-worked and solution-strengthened alloy 5083-H32 (230 MPa, or 33 ksi). The absence of precipitate-forming elements in these low- to moderate-strength non-heat-treatable alloys becomes a positive attribute when considering weldability, because many of the alloy additions needed for precipitation hardening (for example, copper plus magnesium, or magnesium plus silicon) can lead to liquation or hot cracking during welding. In addition, joint efficiencies are higher in non-heat-treatable alloys because the heat-affected zone (HAZ) is not compromised by the coarsening or dissolution of precipitates. This obviates the need for thick joint lands or postweld heat treatment and favors the use of welded structures in the as-welded condition. Selection and Weldability of Non-Heat-Treatable Aluminum Alloys C.E. Cross, Martin Marietta Astronautics Group; M.L. Kohn, FMC Corporation

Alloy Classification and Typical Applications Non-heat-treatable wrought aluminum alloys can be placed into one of three groups using standard Aluminum Association designations:

ALLOY NUMBER 1XXX 3XXX 4XXX 5XXX

ALLOY ADDITION AL (99% MINIMUM PURITY) AL-MN AL-SI AL-MG

Nominal compositions for common non-heat-treatable alloys are given in Table 1.

TABLE 1 COMPOSITION OF WROUGHT NON-HEAT-TREATABLE ALUMINUM ALLOYS ALLOY DESIGNATIO N

COMPOSITION, WT%(A) Si Fe Cu Mn

1050

0.25

1060

0.25

0.4 0 0.3 5

Mg

Cr

Ni

0.05

0.05

0.05

...

...

0.05

0.03

0.03

...

...

Zn

0.0 5 0.0 5

Ti

OTHERS EACH(B TOTAL(C

AL MIN.(D

)

)

)

0.03

0.03(E)

...

99.50

0.03

0.03(E)

...

99.60

1100

3004

0.95 (SI + 0.05FE) 0.20 0.55 (SI + 0.05 FE) 0.15 (SI + 0.10 FE) 0.00 0.0 0.00 6 5 1.00 (SI + 0.05 FE) 0.70 (SI + 0.10 FE) 0.65 (SI + 0.05 FE) 0.03 0.4 0.10 0 0.10 0.4 0.05 0 0.6 0.7 0.050.20 0.30 0.7 0.25

3005

0.6

0.7

0.30

3105

0.6

0.7

0.30

4032

11.0 13.5 4.56.0 9.011.0 11.0 13.0 9.310.7 6.88.2 3.64.6

1.0

0.501.3

0.8

0.30

0.05

0.05

...

...

0.8

0.30

0.05

0.05

...

...

0.8

0.30

0.15

0.10

...

...

0.8

0.15

0.15

0.15

...

0.8

3.34.7 0.25

0.10

...

...

...

0.8

0.10

0.05

0.10 0.30 0.50 -1.1 1.11.8 2.22.8

...

...

0.10

...

0.10

...

0.15 0.35 0.05 0.20

...

1145(G) 1175(H) 1188 1200 1230(H) 1235 1345 1350(J) 3003

4043 4045(L) 4047(L) 4145(L) 4343(L) 4643

0.05

...

...

...

0.05

0.05

...

...

0.02

0.02

...

...

0.01

0.01

...

...

0.05

...

...

...

0.05

0.05

...

...

0.05

0.05

...

...

0.05

0.05

...

...

0.01

...

0.01

...

1.01.5 1.01.5 1.01.5 0.03 -0.8 ...

...

...

...

0.81.3 0.20 -0.6 0.20 -0.8 0.81.3

...

...

0.10

...

0.20

...

0.10

0.50 -1.3

5005

0.30

0.7

0.20

0.20

5050

0.40

0.7

0.20

0.10

5052

0.25

0.4 0

0.10

0.10

5056

0.30

0.4 0

0.10

0.05 0.20

4.55.6

...

...

0.05(F)

0.15

99.00

0.03

0.03(E)

...

99.45

0.02

0.02(I)

...

99.75

0.01

...

...

99.88

0.05

0.05

0.15

99.00

0.03

0.03(E)

...

99.30

0.06

0.03(E)

...

99.35

0.03

0.03(E)

...

99.45

...

0.03(K)

0.10

99.50

...

0.05

0.15

BAL

...

0.05

0.15

BAL

0.10

0.05

0.15

BAL

0.10

0.05

0.15

BAL

...

0.05

0.15

BAL

0.1 0 0.1 0 0.2 0

0.20

0.05(F)

0.15

BAL

0.20

0.05

0.15

BAL

...

0.05(F)

0.15

BAL

0.2 0 0.2 0 0.1 0

...

0.05(F)

0.15

BAL

...

0.05(F)

0.15

BAL

0.15

0.05(F)

0.15

BAL

0.2 5 0.2 5 0.1 0

...

0.05

0.15

BAL

...

0.05

0.15

BAL

...

0.05

0.15

BAL

0.1 0

...

0.05

0.15

BAL

0.1 0 0.0 5 0.0 4 0.0 3 0.1 0 0.1 0 0.1 0 0.0 5 0.0 5 0.1 0 0.2 5 0.2 5 0.4 0 0.2 5

5083

0.40

0.4 0

0.10

0.40 -1.0

4.04.9

5086

0.40

0.5 0

0.10

0.20 -0.7

3.54.5

5154

0.25

0.4 0

0.10

0.10

3.13.9

5183

0.40

0.4 0

0.10

0.50 -1.0

4.35.2

5252

0.08

0.1 0.10 0 0.45 (SI + 0.05 FE)

0.10

2.22.8 3.13.9

5356

0.25

0.4 0

0.10

4.55.5

5454

0.25

0.4 0

0.10

0.05 0.20 0.50 -1.0

5456

0.25

0.4 0

0.10

0.50 -1.0

4.75.5

5457

0.08

0.1 0

0.20

0.81.2

5554

0.25

0.4 0

0.10

0.15 0.45 0.50 -1.0

5556

0.25

0.4 0

0.10

0.50 -1.0

4.75.5

5652

0.40 (SI + 0.04 FE)

0.01

2.22.8

5654

0.45 (SI + 0.05 FE)

0.01

3.13.9

5657

0.08

0.03

0.61.0

5254

0.1 0

0.10

0.01

2.43.0

2.43.0

0.05 0.25 0.05 0.25 0.15 0.35 0.05 0.25 ...

...

0.2 5

0.15

0.05

0.15

BAL

...

0.2 5

0.15

0.05

0.15

BAL

...

0.2 0

0.20

0.05

0.15

BAL

...

0.2 5

0.15

0.05(F)

0.15

BAL

...

...

0.03(E)

0.10

BAL

0.15 0.35 0.05 0.20 0.05 0.20 0.05 0.20 ...

...

0.0 5 0.2 0

0.05

0.05

0.15

BAL

0.06 0.20 0.20

0.05(F)

0.15

BAL

0.05

0.15

BAL

0.05 0.20 0.05 0.20 0.15 0.35 0.15 0.35 ...

...

0.1 0

...

0.2 5

...

0.2 5

0.20

0.05

0.15

BAL

...

0.0 5

...

0.03(E)

0.10

BAL

...

0.2 5

0.05(F)

0.15

BAL

...

0.2 5

0.05(F)

0.15

BAL

...

0.1 0

0.05 0.20 0.05 0.20 ...

0.05

0.15

BAL

...

0.2 0

0.05(F)

0.15

BAL

...

0.0 5

0.02(I)

0.05

BAL

0.05 0.15 ...

Source: Ref 1

(A) COMPOSITION IN MAXIMUM WT% UNLESS SHOWN AS A RANGE OR A MINIMUM. (B) A 0.0008 WT% BE MAXIMUM IS APPLICABLE TO ANY ALLOY TO BE USED AS WELDING ELECTRODE OR WELDING ROD. (C) THE SUM OF THOSE "OTHERS" METALLIC ELEMENTS 0.010% OR MORE EACH, EXPRESSED TO THE SECOND DECIMAL BEFORE DETERMINING THE SUM. (D) THE ALUMINUM CONTENT FOR UNALLOYED ALUMINUM NOT MADE BY A REFINING PROCESS IS THE DIFFERENCE BETWEEN 100.00% AND THE SUM OF ALL OTHER METALLIC

ELEMENTS PRESENT IN AMOUNTS OF 0.010% OR MORE EACH, EXPRESSED TO THE SECOND DECIMAL BEFORE DETERMINING THE SUM. (E) 0.05% V MAXIMUM. (F) 0.0008% BE MAX FOR WELDING ELECTRODE AND WELDING ROD ONLY. (G) FOIL. (H) CLADDING ALLOY. (I) 0.03% GA MAXIMUM; 0.05% V MAXIMUM. (J) ELECTRIC CONDUCTOR (EC). (K) 0.02% (V + TI) MAXIMUM; 0.05% B MAX; 0.03% GA MAXIMUM. (L) BRAZING ALLOY. 1xxx Alloys. The 1xxx-series alloys are of commercial purity (>99% Al) and are used where thermal/electrical conduction or corrosion resistance becomes paramount over strength in design considerations (for example, alloy 1100 is used for sheet metal work, fin stock and chemical equipment). Alloys with purity levels greater than 99.5% are used for electrical conductors (for example, alloy 1350). 3xxx Alloys. The 3xxx-series alloys are used in applications where added additional strength and formability are needed,

in addition to maintaining excellent corrosion resistance (for example, alloy 3004 is used for sheet metal work, storage tanks, and beverage containers). Typical applications include cooking utensils, pressure vessels, and building products (siding, gutters, and so on). These alloys get their strength from cold work and fine (Mn,Fe)Al6 dispersoids that pin grain and subgrain boundaries. There is also a small degree of solid solution strengthening from both manganese and magnesium. 4xxx Alloys. Apart from their use as welding filler material (see the section "Filler Alloy Selection" in this article), the 4xxx-series alloys have limited industrial application in wrought form. 5xxx Alloys. The 5xxx-series alloys are used in cases where still higher strengths are required; this strength is achieved from large quantities of magnesium in solid solution. More importantly, magnesium promotes work hardening by lowering the stacking fault energy, thus reducing the tendency for dynamic recovery. Applications for 5xxx-series alloys include automobile and appliance trim, pressure vessels, armor plate, and components for marine and cryogenic service. The ballistic analysis of alloy 5083 is given in the section "Weld Properties" of this article.

While these alloys normally exhibit good corrosion resistance, care must be taken during processing to avoid formation of continuous β-Mg3Al2 precipitates at grain boundaries, which can lead to intergranular corrosion. This can occur in heavily cold-worked, high-magnesium alloys exposed to temperatures from 120 to 200 °C (250 to 390 °F). Alloy 5454 possesses the highest magnesium content suitable for sustained elevated temperatures and has become the standard alloy used for truck bodies for hot oil or asphalt applications, and storage tanks for heated products. Temper Designations. Appropriate heat treatments and temper designations for non-heat-treatable alloys include a

solution anneal (O temper) and a cold-worked condition (H1 temper). When the cold work achieved during processing is more than desired, a partial anneal can be used to reduce strength (H2 temper). When dealing with 5xxx-series alloys, it is common to use a low-temperature heat treatment (H3 temper) to stabilize the microstructure and prevent intergranular corrosion. This treatment involves heating to 120 °C (250 °F), resulting in partial recovery and making properties less susceptible to change with time. When working with H-type tempers, a second digit is used to indicate the degree of cold work:

RELATIVE HARDNESS TEMPER DESIGNATION FULL HARD H18 3 H16 HARD 4

H14

HARD

H12

HARD

A more thorough and complete description of non-heat-treatable alloys and tempers is given in Volume 2 of the ASM Handbook (Ref 2).


1. ALUMINUM STANDARDS AND DATA, ALUMINUM ASSOCIATION, 1984, P 93 2. PROPERTIES AND SELECTION NONFERROUS ALLOYS AND SPECIAL-PURPOSE MATERIALS, VOL 2, ASM HANDBOOK, ASM INTERNATIONAL, 1990, P 15-28 Selection and Weldability of Non-Heat-Treatable Aluminum Alloys C.E. Cross, Martin Marietta Astronautics Group; M.L. Kohn, FMC Corporation

Filler Alloy Selection Filler alloys used to join non-heat-treatable alloys can be selected from one of three alloy groups:

ALLOY NUMBER 1XXX 4XXX 5XXX

ALLOY ADDITION AL (99% MINIMUM PURITY) A1-SI AL-MG

Commonly used filler alloys include 1100, 1188, 4043, 4047, 5554, 5654, 5183, 5356, and 5556. Nominal compositions for selected filler alloys are included in Table 1. Selecting the best filler alloy for a given application depends on the desired performance relative to weldability, strength, ductility, and corrosion resistance. In general, the filler alloy selected should be similar in composition to the base metal alloy. Thus, a 1xxx filler alloy is recommended for joining 1xxx- or 3xxx-series base metal alloys. Similarly, 5xxx filler alloys are used to join 5xxx-series base metal alloys. An exception to this rule is encountered when weldability becomes an issue. Weldability of non-heat-treatable aluminum alloys can be measured in terms of resistance to hot cracking and porosity. Hot Cracking. Problems with hot cracking are encountered when welding under highly constrained conditions or when

welding certain alloys that are highly susceptible to cracking. Such is the case when welding 5xxx alloys that have a lowrange magnesium content (

1 1 to 2 wt% Mg) where, according to the crack susceptibility curve (Fig. 1), a peak in 2 2

susceptibility occurs. To avoid cracking, use of a high-magnesium filler alloy is recommended.

FIG. 1 PLOT OF HOT CRACKING SENSITIVITY VERSUS MAGNESIUM CONTENT TO SHOW PEAK IN SUSCEPTIBILITY OBTAINED FOR ALUMINUM-MAGNESIUM (5XXX) ALLOYS. SOURCE: REF 3

Similar problems may be encountered when 1xxx fillers are used to join 5xxx alloys (or vice versa), or when welding dissimilar metal alloys such as alloys 1100 and 5083, where mutual dilution may result in low magnesium levels. Electron-beam welding or laser-beam welding can also result in cracking when magnesium, a high-vapor-pressure alloying element, is boiled off. The problem is aggravated when welding in a vacuum environment. An example of a center-line hot crack produced when electron-beam welding alloy 5083 is shown in Fig. 2, where magnesium levels were measured to be approximately 3 wt% in the vicinity of the crack, based on microprobe analysis. This was an autogenous weld made in a hard vacuum (0.13 mPa, or 10-6 torr).

FIG. 2 CENTERLINE HOT CRACK PRODUCED IN AUTOGENOUS ELECTRON-BEAM WELDED ALLOY 5083

Another approach to be taken when hot cracking persists is to use 4xxx fillers. These aluminum-silicon alloys have exceptional resistance to cracking, due in part to their abundance of liquid eutectic available for back-filling (Ref 4). However, their use should be avoided when welding high-magnesium alloys (>3 wt%) because of embrittlement from excessive Mg2Si precipitation. Other drawbacks include low joint ductility and nonmatching color when anodized. Weldability will improve with increased silicon content (for example, alloy 4047 versus alloy 4043), but the disadvantages discussed above will likewise become more predominant. Porosity. Non-heat-treatable aluminum alloys are susceptible to hydrogen-induced weld metal porosity, as are all

aluminum alloys in general. This porosity forms during solidification due to the abrupt drop in hydrogen solubility when going from liquid to solid. Porosity can best be avoided by minimizing hydrogen pickup during welding. This can be accomplished through proper joint preparation (that is, removal of hydrocarbons), use of high-grade (low-dew-point) shielding gas, and careful storage of filler wire (that is, protection from exposure to moisture and oil). It has been determined that welding filler wire is often the primary source of hydrogen contamination. The 5xxx-series filler alloys, in particular, are susceptible to the hydration of surface oxides, which can result in porosity (Ref 5).

In summary, guidelines for filler alloy selection when welding non-heat-treatable alloys are:

BASE METAL 1XXX 3XXX 5XXX (LOW MAGNESIUM) 5XXX (HIGH MAGNESIUM)

FILLER METAL 1XXX, 4XXX 1XXX, 4XXX 5XXX, 4XXX 5XXX


3. J.H. DUDAS AND F.R. COLLINS, PREVENTING WELD CRACKS IN HIGH-STRENGTH ALUMINUM ALLOYS, WELD. J., VOL 45, 1966, P 241S-249S 4. C.E. CROSS, L.S. KRAMER, W.T. TACK, AND L.W. LOECHEL, ALUMINUM WELDABILITY AND HOT TEARING THEORY, WELDABILITY OF MATERIALS, ASM INTERNATIONAL, 1990, P 275-282 5. R.P. MARTUKANITZ AND P.R. MICHNUK, SOURCES OF POROSITY IN GAS METAL ARC WELDING OF ALUMINUM, TRENDS IN WELDING RESEARCH, ASM INTERNATIONAL, 1982, P 315-330 Selection and Weldability of Non-Heat-Treatable Aluminum Alloys C.E. Cross, Martin Marietta Astronautics Group; M.L. Kohn, FMC Corporation

Weld Properties When non-heat-treatable alloys are welded, microstructural damage is incurred in the HAZ. Unlike the case of heattreatable alloys, whose strengthening precipitates may dissolve or coarsen, the HAZ damage in non-heat-treatable alloys is limited to recovery, recrystallization, and grain growth. Thus, loss in strength in the HAZ is not nearly as severe as that experienced in heat-treatable alloys. (An example of this is shown in the hardness traverses of Fig. 8 in the article "Selection and Weldability of Heat-Treatable Aluminum Alloys" in this Volume, where alloy 5456 is compared against alloys 6061 and 2219.) For this reason, 5xxx-series alloys are popular for use in welded pressure vessels where reasonable joint strengths can be obtained in the as-welded condition without the need for post-weld heat treatment. The weld metal of non-heat-treatable aluminum alloys is typically the weakest part of the joint and is the location of failure when the joint is loaded in tension. This is in contrast to most heat-treatable aluminum alloys, where the heataffected zone often is the weakest link. The weld metal microstructure of the non-heat-treatable alloys consists of columnar, epitaxial grains with a cellular or columnar-dendritic substructure that has interdendritic eutectic constituentsprimarily (Fe,Mn)Al6 for 1xxx and 3xxx alloys; Si for 4xxx alloys; and Mg3Al2 for 5xxx alloys. Joint efficiencies for several non-heat-treatable alloys are compared in Table 2. It should be noted that joint efficiency is highest for annealed plate (approaching 100%) and decreases continuously with the degree of cold working. Also shown in Table 2 is the effect of composition and cold work on plate properties. The two 5xxx-series alloys shown exhibit significant improvement in strength over 1xxx- and 3xxx-series alloys, a reflection on the strengthening ability of magnesium. Likewise, 5xxx-series alloys with higher magnesium content have higher strength (for example, alloy 5083 versus alloy 5050). This same trend is observed for weld metal strength. For example, a weld made on high-magnesium alloy 5083 (4% Mg) is much stronger than a similar weld made on alloy 5050 (1% Mg).

TABLE 2 EFFECT OF VARIATIONS IN TEMPER AND FILLER ALLOY ON NON-HEAT-TREATABLE ALUMINUM GAS-METAL ARC WELDMENT PROPERTIES

BASE ALLOY AND TEMPE R

PLATE PROPERTIES YIELD ULTIMAT STRENGT E H TENSILE STRENGT H MPA KSI MPA KSI

1100-O 34

5

90

13

ELONGATIO N IN 50 MM (2 IN.), %

FILLE R ALLO Y

WELD PROPERTIES YIELD ULTIMATE TENSILE STRENGT STRENGTH H

ELONGATIO N IN 50 MM (2 IN.), %

JOINT EFFICIENC Y, %(A)

MPA

KSI

MP A

KSI

13. 5 13 13 13. 2 16

23

100

21 20 17

100 81 83

15

55

21 23 22 23. 5 41 40 43. 5

20 14.5 15 14.5

100 92 88 73

17 16 12

98 87 95

35

1100

41

6

93

41 41 48

6 6 7

90 90 91

3003-O 41

6

110

16

30

4043 1100 4043

3003H18 5050-O 5050H32 5050H38 5083-O 5083H32

186

27

200

29

4

1100

59

8.5

110

55 145

8 21

145 172

21 25

24 9

200

29

221

32

6

5356 5356 4043 5356

55 97 90 97

8 14 13 14

145 159 152 162

145 228

21 33

290 317

42 46

22 16

5356 5356 5183

145 145 152

21 21 22

283 276 300

Source: Ref 6

(A) JOINT EFFICIENCY = (UTS OF WELD)/(UTS OF PLATE), WHERE UTS STANDS FOR ULTIMATE TENSILE STRENGTH. The correspondence between plate and joint strength does not apply to temper. For the high-magnesium alloy 5083 plate, no increase in joint strength is observed when welding an annealed versus a cold-worked temper (see Table 2). Likewise, filler metal selection has little effect on joint strength, as shown in Table 2. However, it is clear that use of a 4xxx filler does result in ductility loss. Also, filler alloys high in iron, silicon, or manganese (for example, filler alloy 5183) will have greater amounts of eutectic constituents, which will result in lower weld ductility. Even so, non-heat-treatable alloy welds exhibit significantly better ductility in comparison to similar heat-treatable welds. Example 1: Use of Welded Alloy 5083 in Ballistic Applications. An important application for heavy-gage alloy 5083 is the construction of tactical military vehicles. The hulls and turrets of Army vehicles such as the M113 armored personnel carrier, the M2/M3 infantry and cavalry fighting vehicles, the M109 self-propelled howitzer, and the Marine Corps AAV7A amphibians all consist of welded 5083 aluminum structures. There are also a myriad of brackets, clips, and so on, welded to the hulls and turrets, although not normally fabricated to ballistic requirements. Aluminum alloys considered for use as armor include alloys 5083, 7039, and 2519. Alloys 7039 and 2519 are weldable and heat-treatable, capable of yield strengths exceeding 330 MPa (48 ksi). However, alloy 2519 weldments have failed rigorous ballistic shock tests (MIL-STD-1946, discussed below), even though the base material exhibits excellent ballistic and mechanical properties. Alloy 7039 possesses superior ballistic properties compared to alloy 5083, but it suffers from a tendency to delaminate and to exhibit stress-corrosion cracking. Alloy 5083 weldments have passed ballistic weld tests, and they possess the necessary resistance to corrosion. In the MIL-STD-1946 ballistic test, a weldment is subjected to impact by a blunt-nose 75 mm (3 in.) aluminum projectile, fired at a velocity specified according to material thickness and alloy composition. Butt-welded specimens are struck no more than 13 mm (

1 in.) from the weld centerline. Corner joints are struck so that the center of the impact projectile is 2

within 50 ± 13 mm (2 ±

1 in.) from the weld toe on its back side. Welds are considered acceptable when the total length 2

of cracking generated is less than 305 mm (12 in.). An offset double-V configuration of the butt joint is commonly used in ballistic applications (Fig. 3). This results in solid metal backing for both welds and allows for a slip fit between adjoining plates, thereby simplifying the fit-up of large, heavy plates.

FIG. 3 OFFSET DOUBLE-V JOINT CONFIGURATION USED TO WELD THICK SECTION ALUMINUM ARMOR PLATE


6. "WELDING KAISER ALUMINUM," KAISER ALUMINUM & SALES, INC., OAKLAND, CA, 1978, P 3.25-3.26 Selection and Weldability of Non-Heat-Treatable Aluminum Alloys C.E. Cross, Martin Marietta Astronautics Group; M.L. Kohn, FMC Corporation

References

1. ALUMINUM STANDARDS AND DATA, ALUMINUM ASSOCIATION, 1984, P 93 2. PROPERTIES AND SELECTION NONFERROUS ALLOYS AND SPECIAL-PURPOSE MATERIALS, VOL 2, ASM HANDBOOK, ASM INTERNATIONAL, 1990, P 15-28 3. J.H. DUDAS AND F.R. COLLINS, PREVENTING WELD CRACKS IN HIGH-STRENGTH ALUMINUM ALLOYS, WELD. J., VOL 45, 1966, P 241S-249S 4. C.E. CROSS, L.S. KRAMER, W.T. TACK, AND L.W. LOECHEL, ALUMINUM WELDABILITY AND HOT TEARING THEORY, WELDABILITY OF MATERIALS, ASM INTERNATIONAL, 1990, P 275-282 5. R.P. MARTUKANITZ AND P.R. MICHNUK, SOURCES OF POROSITY IN GAS METAL ARC WELDING OF ALUMINUM, TRENDS IN WELDING RESEARCH, ASM INTERNATIONAL, 1982, P 315-330 6. "WELDING KAISER ALUMINUM," KAISER ALUMINUM & SALES, INC., OAKLAND, CA, 1978, P 3.25-3.26 Selection and Weldability of Dispersion-Strengthened Aluminum Alloys K. Sampath, Concurrent Technologies Corporation; W.A. Baeslack III, The Ohio State University

Introduction CONVENTIONAL HIGH-STRENGTH ALUMINUM ALLOYS produced via ingot metallurgy (I/M) processing contain precipitate particles that interact with moving dislocations, thus providing high strength at room temperature. However, when these alloys are exposed to elevated temperatures, the precipitate particles undergo dissolution and/or coarsening,

thereby reducing the strength of the base alloy. This behavior precludes the use of these alloys for elevated-temperature applications. In recent years, advances in two broad powder metallurgy (P/M) technologies--rapid solidification (RS) and mechanical alloying (mechanical attrition)--have enabled the development of a new family of high-strength aluminum alloys for elevated-temperature applications (Ref 1). These new high-strength P/M aluminum alloys depend primarily on dispersion strengthening. Unlike the precipitate particles in I/M aluminum alloys, the dispersoid particles resist dissolution and/or coarsening when the alloys are exposed to elevated temperatures. In addition to high strength, these dispersion-strengthened aluminum alloys exhibit light weight, appreciable ductility, excellent fracture toughness, and resistance to stress-corrosion cracking (Ref 2). These alloys also exhibit coefficients of thermal expansion and galvanic properties comparable to those of conventional structural aluminum alloys (Ref 2). In view of the above properties, these dispersion-strengthened aluminum alloys are considered candidate materials to replace conventional high-strength aluminum alloys and titanium alloys used in the manufacture of selected aerospace components such as fan and compressor cases; vanes and blades in gas turbine engines; and fins, winglets, and rocket motor cases of missiles (Ref 3). Use of these alloys is expected to increase the thrust-to-weight ratio of gas turbine engines and allow significant improvements in the performance characteristics of advanced aircraft and weapons systems. High-strength aluminum alloys produced via RS-P/M technology am primarily based on hypereutectic aluminum-iron compositions, although other alloys based on aluminum-chromium-zirconium, aluminum-manganese, aluminumtitanium, and aluminum-beryllium systems are also under various stages of development and/or evaluation (Ref 1). The dispersion-strengthened aluminum alloys produced via mechanical alloying contain a fine distribution of oxides, oxynitrides, and/or carbides in a fine-grain α-aluminum matrix. The α-aluminum matrix may additionally be strengthened by solid-solution strengthening or work hardening (Ref 1). Although many dispersion-strengthened aluminum alloys produced via mechanical alloying do not exhibit strength comparable to that of the conventional high-strength I/M aluminum alloys at room temperature, these alloys exhibit a minimal decrease in strength with increasing temperature (Ref 1). The twofold purpose of this article is: (1) to guide users in the selection and application of welding processes and conditions for the joining of dispersion-strengthened aluminum alloys, based primarily on an understanding of the weldability of the material; and (2) to provide specific examples of material responses to welding conditions, chiefly by highlighting the microstructural development in the weld zone. In view of the limited number of publications available on the weldability and metallurgical response of this broad class of alloys, this article refers exclusively to the metallurgy and weldability of alloys based on the aluminum-iron system that are produced using various RS-P/M processing techniques.

References

1. J.R. PICKENS, HIGH STRENGTH ALUMINUM P/M ALLOYS, 10TH ED., VOL 2, METALS HANDBOOK, PROPERTIES AND SELECTION: NON FERROUS ALLOYS AND SPECIAL-PURPOSE MATERIALS, ASM INTERNATIONAL, 1990, P 200-215 2. S.K. DAS AND L.A. DAVIES, HIGH PERFORMANCE AEROSPACE ALLOYS VIA RAPID SOLIDIFICATION PROCESSING, MATER. SCI. ENG., VOL 98, 1988, P 1-12 3. C.M. ADAM, RAPID SOLIDIFICATION PROCESSING METHODS, MECHANICAL BEHAVIOR OF RAPIDLY SOLIDIFIED MATERIALS, S.M.L. SASTRY AND B.A. MACDONALD, ED., THE METALLURGICAL SOCIETY OF AIME, 1985, P 21-39 Selection and Weldability of Dispersion-Strengthened Aluminum Alloys K. Sampath, Concurrent Technologies Corporation; W.A. Baeslack III, The Ohio State University

Metallurgy of RS-P/M Aluminum-Iron-Base Alloys Table 1 shows a list of selected dispersion-strengthened aluminum-iron-base alloys produced via RS-P/M processing. These alloys contain ternary and quaternary additions of transition and rare earth metals such as molybdenum, nickel,

cerium, vanadium, and silicon. These alloy additions exhibit low solid-solubility and low solid-state diffusivity in aluminum (Ref 4), and consequently they tend to form thermally stable dispersoid particles. However, these alloy compositions are characterized by a wider freezing range (typically in excess of 300 °C, or 540 °F) than most aluminum alloys. Conventional I/M processing of these compositions produces equilibrium or near-equilibrium microstructures containing coarse Al3Fe-type primary intermetallic particles in a coarse -aluminum matrix. Such microstructures exhibit poor ductility and cannot be used for structural engineering applications. In contrast, RS processing of these novel chemistries suppresses the formation of Al3Fe-type primary intermetallic particles. It instead produces particulates exhibiting refined metastable microstructures, such as an Al-Al6Fe-type microeutectic and/or a solute-supersaturated dendritic aluminum. In other words, RS processing conditions produce refined hypoeutectic-type microstructures from hypereutectic melts. The occurrence of these metastable hypoeutectic-type microstructures is attributed to the following two characteristics of RS processing techniques: (1) subdivision of the melt into particulates isolates potential nucleants and produces nucleant-free particulates that undergo nonequilibrium solidification (Ref 5); and (2) the rapid cooling rate does not provide sufficient reaction time for a potential nucleant to produce the equilibrium microstructure (Ref 6).

TABLE 1 CHEMICAL COMPOSITIONS OF DISPERSION-STRENGTHENED ALUMINUM-BASE ALLOYS PRODUCED BY RS-P/M

NOMINAL CHEMICAL COMPOSITION, WT% AL-8FE-1.7NI AL-8.4FE-3.7CE AL-9FE-4CE (AA8019 ALLOY) AL-9FE-7CE AL-10FE-5CE AL-8.5FE-1.3V-1.7SI (AA8009 ALLOY) AL-11.7FE-1.2V-2.4SI (FVS 1212 ALLOY) AL-8FE-2.3MO AL-8.7FE-2.8MO-1V

MANUFACTURER ALCOA ALCOA ALCOA ALCOA ALCOA ALLIEDSIGNAL INC. ALLIEDSIGNAL INC. PRATT & WHITNEY PRATT & WHITNEY

Subsequent thermomechanical processing of the rapidly solidified particulates produces a consolidated product with a unique "engineered" microstructure consisting of submicron-size dispersoid particles in an extremely fine-grain (2 >0.5-2

... ...

... ...

... ...

... ...

... ...

10B

VACUUM

...

...

...

...

...

VACUUM

...

...

>0.001- . . . 0.5 ≤ 0.001 . . .

...

10C

>265 >65265 >0.1365 ≤ 0.13

...

...

...

...

(A)

RT, ROOM TEMPERATURE

TABLE 4 CLASSIFICATION OF SOLDERING FLUXES FLUX TYPE

USES

CORROSIVE (INORGANIC) ACIDS STRUCTURAL

SALTS

STRUCTURAL

INTERMEDIATE (ORGANIC) ACIDS STRUCTURAL, ELECTRICAL HALOGENS

STRUCTURAL, ELECTRICAL

AMINES AND AMIDES

STRUCTURAL, ELECTRICAL

NONCORROSIVE (ORGANIC) ROSIN ELECTRICAL NONACTIVATED (R) ROSIN MILDLY ELECTRICAL ACTIVATED (RMA) ROSIN ACTIVATED (RA)

ELECTRICAL, STRUCTURAL

CONSTITUENTS

TEMPERATURE STABILITY

RECOMMENDED TEMPERATURE RANGE °C °F

HYDROCHLORIC, HYDROFLUORIC, ORTHOPHOSPHORIC ZINC CHLORIDE, AMMONIUM CHLORIDE, TIN CHLORIDE

GOOD

150450

300840

EXCELLENT

150450

300840

LACTIC, OLEIC, STEARIC, GLUTAMIC, PHTHALIC ANILINE HYDROCHLORIDE, GLUTAMIC HYDROCHLORIDE, BROMIDE DERIVATIVES OF PALMITIC ACID, HYDRAZINE HYDROCHLORIDE (OR HYDROBROMIDE) UREA, ETHYLENE DIAMINE, MONOTRIETHANOL AMINE, TRIETHANOLAMINE

FAIRLY GOOD

180280

355500

FAIRLY GOOD

180280

355500

FAIR

180280

355500

WATER-WHITE ROSIN

POOR

140270

285520

ROSIN WITH ACTIVATOR (HALIDEBASE, ORGANIC ACID, AMINES, AMIDES . . .) ROSIN WITH MORE ACTIVATOR THAN RMA FLUX

POOR

140270

285520

FAIR

140270

285520

Source: Ref 5 Additional Metallurgical Considerations. Exposure to brazing temperatures can also result in base metal embrittlement. These embrittling conditions must be taken into consideration in the design of a brazed joint and brazing process parameters. Soldering occurs at lower temperatures than brazing, so embrittlement seldom happens in the base metal of a soldered joint because it is thermally activated. Embrittlement in the base metal can be caused by second-phase precipitation, hydrogen embrittlement, or the brazing of dissimilar metals. Second-Phase Precipitation. The most common embrittlement phenomenon is second-phase precipitation of carbides

at grain boundaries. Plain steels with low carbon contents can be embrittled by the precipitation of a cementite (Fe3C) film at the grain boundary. For a steel with 0.02 wt% C, this phenomenon occurs at approximately 400 °C (750 °F). Ferritic and austenitic stainless steels can have their corrosion behavior and toughness impaired by chromium carbide precipitation between 450 to 850 °C (750 to 1560 °F). This precipitation forms a chromium-depleted zone around the carbides and along the grain boundaries, which decreases the corrosion resistance of the alloy. This phenomenon is called sensitization. For the austenitic stainless steels, a low-carbon steel or a stabilized steel should be chosen to avoid the loss in corrosion resistance. Higher chromium diffusivity in the ferritic stainless steels results in a narrower chromiumdepleted zone than that observed in the austenitic stainless steels. Thus, a heat treatment can remove the chromiumdepleted zone and restore corrosion resistance. However, the toughness of the ferritic stainless steels will not be recovered by heat treatment. Aluminum-magnesium alloys can experience an Al3Mg2 grain boundary precipitate that decreases the corrosion resistance of the alloy. In this case the difference between the corrosion potentials of the precipitate and the matrix increases the corrosion rate at the matrix-precipitate interface. Depending on the conditions, intergranular corrosion may take place, followed by stress-corrosion cracking. Hydrogen Embrittlement. Hydrogen can cause severe embrittlement in many different materials. It can diffuse faster at

room temperature and may cause a delayed embrittlement, mainly in steels. This type of embrittlement is not important for soldering. In steels, hydrogen embrittlement occurs when four conditions are met: favorable microstructure (martensite, bainite, ferrite), hydrogen level, residual stress, and low temperature (50 TO 100 ≤ 50

>2 TO 4

38

93

OPTIONAL

≤2

AMBIENT(B)

AMBIENT(B)

OPTIONAL

>50 TO 100 ≤ 50

>2 TO 4

38

121

250

OPTIONAL

≤2

AMBIENT(B)

38

100

OPTIONAL

>2 TO 4

93

149

300

OPTIONAL

0.18-0.25

>50 TO 100 ≤ 25

≤1

AMBIENT(B)

38

100

OPTIONAL

>1 TO 2 >2 TO 4

38

100

93

200

OPTIONAL

93

200

149

300

OPTIONAL

0.25-0.29

>25 TO 50 >50 TO 100 ≤ 13

10

50

65

150

OPTIONAL

WITHIN SPECIFICATION


≤

1 2

100

200

150

TEMPERATURE RANGE °C °F

>2 TO 4

100

65

200

PEENING MAY BE NECESSARY

DESIRABLE OR OPTIONAL

>50 TO 100 ≤ 50


100

POSTWELD HEAT TREATMENT

OPTIONAL

YES

YES

YES

YES

... YES

>13 TO 25

1025

1026, 1029, 1030

>25 TO 50 >50 TO 100 ≤ 25

1 2

TO 1 >1 TO 2 >2 TO 4

65

150

149

300

OPTIONAL

590675

11001250

...

121

250

149

300

OPTIONAL

300

177

350

OPTIONAL

11001250 11001250

...

149

590675 590675 590675 590675 590675

11001250 11001250 11001250

...

YES

≤1

AMBIENT(B)

AMBIENT(B)

OPTIONAL

>1 TO 2 >2 TO 4

AMBIENT(B)

38

100

OPTIONAL

38

93

200

OPTIONAL

0.25-0.30

>25 TO 50 >50 TO 100 ≤ 25

≤1

AMBIENT(B)

38

100

>1 TO 2 >2 TO 4

38

100

93

200

93

200

149

300

11001250 11001250 11001250

...

0.31-0.34

>25 TO 50 >50 TO 100 ≤ 13

DESIRABLE 590675 DESIRABLE 590675 DESIRABLE 590675

1 2 1 > 2

AMBIENT(B)

38

100

100

93

200

11001250 11001250

...

38

DESIRABLE 590675 DESIRABLE 590675

93

200

149

300

250

177

350

11001250 11001250

YES

121


AMBIENT(B)

38

100

93

200

11001250 11001250

...

38



>13 TO 25

1035, 1037

>


>25 TO 50 >50 TO 100 ≤ 13 >13 TO 25

≤

TO 1 >1 TO 2 >2 TO 4 1 2

≤

>

1 2

TO 1

100

100

... YES

YES YES

...

YES

...

1036, 1041

0.30-0.35

>25 TO 50 >50 TO 100 ≤ 13 >13 TO 25

0.36-0.40

>25 TO 50 >50 TO 100 ≤ 13 >13 TO 25

0.41-0.44

>25 TO 50 >50 TO 100 ≤ 13 >13 TO 25

1038, 1039, 1040

0.34-0.40

>25 TO 50 >50 TO 100 ≤ 13

>1 TO 2 >2 TO 4

93

200

149

300

149

300

205

400

1 2 1 > 2

38

100

93

200

65

150

149

300

121

250

149

300

149

300

177

350

65

150

93

200

93

200

149

300

149

300

177

350

177

350

205

400

93

200

149

300

149

300

177

350

177

350

205

400

205

400

230

450

38

100

93

200

≤

TO 1 >1 TO 2 >2 TO 4 1 2

≤

>

1 2

TO 1 >1 TO 2 >2 TO 4 1 2

≤

>

1 2

TO 1 >1 TO 2 >2 TO 4 ≤

1 2


11001250 11001250

YES


11001250 11001250

...


1001250 11001250

YES


11001250 11001250

...


1001250 11001250

YES


11001250 11001250

...


11001250 11001250

YES

DESIRABLE 590675

11001250

...

YES

...

YES

...

YES

...

YES

>13 TO 25

0.41-0.44

>25 TO 50 >50 TO 100 ≤ 13 >13 TO 25

1042, 1043


>25 TO 50 >50 TO 100 ≤ 13 >13 TO 25

1044, 1045, 1046


1048, 1049, 1050, 1052, 0.43-0.50 1053

>25 TO 50 >50 TO 100 ≤ 13

1 2

>

TO 1 >1 TO 2 >2 TO 4 1 2

≤

>

1 2

TO 1 >1 TO 2 >2 TO 4 1 2

≤

>

1 2

TO 1 >1 TO 2 >2 TO 4 1 2

≤

>13 TO 100 ≤ 25

>

>25 TO 50

>1 TO 2

1 2

TO 4 ≤1

93

200

149

300

DESIRABLE 590675

11001250

...

121

250

177

350

300

205

400

11001250 11001250

YES

149


65

150

121

250

200

149

300

11001250 11001250

...

93


149

300

177

350

400

230

450

11001250 11001250

YES

205


93

200

149

300

250

149

300

11001250 11001250

YES

121


149

300

177

350

400

230

450

11001250 11001250

YES

205


149

300

177

350

400

230

450

11001250 11001250

YES

205


93

200

149

300

300

177

350

11001250 11001250

YES

149


YES

...

YES

YES

YES

YES

YES

1108, 1109, 1110

1116, 1117, 1118, 1119



1132, 1137, 1139, 1140, 0.27-0.30 1141, 1144, 1145, 1146, 1151

0.31-0.35

>50 TO 100 ≤ 50

>2 TO 4

205

≤2

AMBIENT(B)

>50 TO 100 ≤ 25

>2 TO 4

38

≤1

AMBIENT(B)

>25 TO 100 ≤ 13

>1 TO 4

93

200

10

50

>13 TO 25

>

38

100

93

200

121

250

38

100

65

150

121

250

149

300

65

150

93

200

>25 TO 50 >50 TO 100 ≤ 13 >13 TO 25

0.36-0.40

>25 TO 50 >50 TO 100 ≤ 13 >13 TO 25

1 2

≤ 1 2

TO 1 >1 TO 2 >2 TO 4 1 2

≤

>

1 2

TO 1 >1 TO 2 >2 TO 4 1 2

≤

>

1 2

400

100

230

450

DESIRABLE 590675

11001250

YES

NOT RECOMMENDED NOT RECOMMENDED

OPTIONAL

590675 590675

11001250 11001250

...


OPTIONAL

590675 590675

11001250 11001250

...


OPTIONAL

590675 DESIRABLE 590675

11001250 11001250

...



11001250 11001250

YES


DESIRABLE 590675 OPTIONAL 590675

11001250 11001250

...



11001250 11001250

YES



11001250 11001250

...

OPTIONAL

OPTIONAL

YES

YES

...

YES

...

YES

...

0.41-0.45

>25 TO 100 ≤ 13

149

300

NOT RECOMMENDED

DESIRABLE 590675

11001250

YES

93

200 300


11001250 11001250

...

149


177

350

NOT RECOMMENDED

DESIRABLE 590675

11001250

YES

149

300 400


11001250 11001250

...

205


TO 1 >1 TO 4

230

450

NOT RECOMMENDED

OPTIONAL

590675

11001250

YES

≤2

AMBIENT(B)

OPTIONAL

>2 TO 4

38

590675 590675

11001250 11001250

...

>50 TO 100 ≤ 50


≤2

AMBIENT(B)

OPTIONAL

>2 TO 4

38

100

590675 590675

11001250 11001250

...

>50 TO 100 ≤ 13


121

250

>

149

300


10251050 10251050

...

>13 TO 25


177

350

NOT RECOMMENDED

DESIRABLE 550565

10251050

YES

>13 TO 25

0.45-0.50

>25 TO 100 ≤ 13 >13 TO 25

1211, 1212, 1213, 1215, WITHIN B1111, B1112, B1113 SPECIFICATION

12L 13(C), 12L 14(C)

1330, 1335, 1340, 1345


0.27-0.33

TO 1 >1 TO 4

>25 TO 100 ≤ 50

>25 TO 50

1 2

≤

>

1 2

TO 1 >1 TO 4 1 2

≤

>

1 2

1 2

≤ 1 2

TO 1 >1 TO 2

100

OPTIONAL

OPTIONAL

...

...

YES

YES

...

0.33-0.38

≤ 13

>13 TO 25

0.38-0.43

>25 TO 50 ≤ 13 >13 TO 25

0.43-0.49

>25 TO 50 ≤ 13 >13 TO 25

1513, 1518, 1522, 1525

1 2

≤

>

1 2

TO 1 >1 TO 2 1 2

≤

>

1 2

TO 1 >1 TO 2 1 2

≤

>

1 2

TO I >1 TO 2 ≤1

149

300



10251050 10251050

...

205

400

205

400

NOT RECOMMENDED NOT RECOMMENDED NOT RECOMMENDED


10251050 10251050 10251050

YES

177

350

230

450

290(D) 550(D) NOT RECOMMENDED 205 400 NOT RECOMMENDED 260(D) 500 NOT RECOMMENDED


10251050 1025 1050 10251050

YES

600(D) NOT RECOMMENDED (B) AMBIENT AMBIENT(B)

DESIRABLE 550565 OPTIONAL 590675 OPTIONAL 590675 OPTIONAL 590675

10251050 11001250 11001250 11001250

YES

590675 590675 590675

11001250 11001250 11001250

...

590675

11001250

...

≤ 0.20

>25 TO 50 ≤ 25

> 1 AMBIENT(B) TO 2 >2 38 100 TO 4

38

100

93

200

0.21-0.25

>25 TO 50 >50 TO 100 ≤ 25

≤1

AMBIENT(B)

38

100

OPTIONAL

>1 TO 2 >2 TO 4

38

100

93

200

OPTIONAL

93

200

149

300

OPTIONAL

0.26-0.29

>25 TO 50 >50 TO 100 ≤ 13

10

50

65

150

OPTIONAL

≤

1 2

315

...

... ...

... ...

... ... YES

... YES

>13 TO 25

1524, 1526, 1527

1536, 1541

65

150

149

300

OPTIONAL

590675

11001250

...

121

250

149

300

OPTIONAL

300

177

350

OPTIONAL

11001250 11001250

...

149

590675 590675

≤1

10

50

38

100

OPTIONAL

590675 590675 590675

11001250 11001250 11001250

...

>

1 2

TO 1 >1 TO 2 >2 TO 4

≤ 0.25

>25 TO 50 >50 TO 100 ≤ 25

>1 TO 2 >2 TO 4

38

100

93

200

OPTIONAL

93

200

149

300

OPTIONAL

0.26-0.29

>25 TO 50 >50 TO 100 ≤ 25

≤1

65

150

121

250

>1 TO 2 >2 TO 4

121

250

149

300

149

300

177

350

11001250 11001250 11001250

...

≤ 0.35

>25 TO 50 >50 TO 100 ≤ 13


38

100

93

200

>

65

150

149

300

11001250 11001250

...

>13 TO 25


121

250

149

300

300

177

350

11001250 11001250

YES

149


65

150

93

200

200

149

300

11001250 11001250

...

93


0.364-0.40

>25 TO 50 >50 TO 100 ≤ 13 >13 TO 25

1 2

≤ 1 2

TO 1 >1 TO 2 >2 TO 4 1 2

≤

>

1 2

TO 1

YES

... YES

... YES

...

YES

...

0.414-0.44

1547, 1548

0.43-0.52

>25 TO 50 >50 TO 100 ≤ 13

>1 TO 2 >2 TO 4

>13 TO 25

>

>25 TO 50 >50 TO 100 ≤ 13 >13 TO 100

1 2

≤ 1 2

TO 1 >1 TO 2 >2 TO 4 1 2

≤

>

1 2

149

300

177

350

177

350

205

400

93

200

149

300

149

300

177

350

177

350

205

400

205

400

230

450

149

300

177

350

205

400

230

450


11001250 11001250

YES


11001250 11001250

...


11001250 11001250

YES


11001250 11001250

...

YES

...

YES

YES

TO 4

Source: Ref 5

(A) RANGES ARE INCLUSIVE. (B) AMBIENT ABOVE -12 °C (10 °F). (C) DUE TO LEAD CONTENT, MANUFACTURING OPERATIONS INVOLVING ELEVATED TEMPERATURES IN THE RANGE OF THOSE ENCOUNTERED IN GAS CUTTING OR WELDING SHOULD BE CARRIED OUT WITH ADEQUATE VENTILATION. (D) HOLD AT TEMPERATURE FOR 1 H AFTER WELDING IS COMPLETED.

Particular care must be exercised when dealing with weldments made at low heat inputs (high cooling rates). Tack welds are usually short, low-heat-input welds that cool rapidly. The first pass of a multipass weld in a large structure may be made at low heat input (to avoid bum-through) or on a relatively cold workpiece. Again, cooling rates may be high. It is essential that proper preheat be used under these circumstances, because the likelihood of martensite formation is much greater, given the rapid cooling rates. Heat input control can also be used to reduce the cooling rate. The use of a higher heat input to slow cooling may not

always be commensurate with good welding practice. For example, the resultant bead size may be too large, or the welding parameters (current, voltage, travel speed) needed to achieve the required heat input level may be outside the operating window of the electrode. Postweld heat treatment may involve holding the completed weld at or near the interpass temperature for a period of

time sufficient to allow hydrogen to escape the weld area. This temperature should be above the temperature range where cracking occurs (approximately 200 °C, or 390 °F). If the interpass and postweld soak temperature are held above the martensite start temperature, Ms, for a sufficient length of time, isothermal transformation to a less susceptible microstructure will occur (for example, bainitic structure with some ferrite). A tempered martensite, however, will generally have better toughness than the isothermal transformation product, so this procedure should be used only when a stress relief is not possible. The use of a proper postweld treatment may allow a reduction in the preheat temperature because both treatments serve the same purpose (that is, reducing the weld metal cooling rate). A lower preheat will increase operator comfort. In the event that martensite formation cannot be avoided (as may be the case in a medium- to high-carbon steel), or when joint restraint is high, it is advisable to postweld heat treat the weldment to produce a less brittle, tempered martensite. The weldment should be held at or near the interpass temperature until postweld heat treatment can be done. This will reduce the possibility of cracking before the postweld heat treatment. Postweld heat treating of carbon steels is normally performed in the 600 to 650 °C (1110 to 1200 °F) range. The soak time is usually 1 h/in. of the thickest member. Many codes (for example, the ASME Boiler and Pressure Vessel Code) define when postweld heat treatment is to be done. Figure 4 can also be used as a guide for determining when postweld heat treatment should be done.

FIG. 4 PLOT OF BASE-METAL CARBON CONTENT VERSUS BASE-METAL THICKNESS OF CARBON STEELS TO SHOW DEGREE OF STRESS RELIEF REQUIRED. A, POSTWELD HEATING IS SELDOM REQUIRED. B, POSTWELD HEATING IS REQUIRED ONLY FOR DIMENSIONAL STABILITY, AS WHEN PARTS ARE TO BE FINISH MACHINED AFTER WELDING. C, POSTWELD STRESS RELIEF IS HIGHLY DESIRABLE FOR REPETITIVELY LOADED OR SHOCK-LOADED STRUCTURES AND FOR RESTRAINED JOINTS HAVING A THICKNESS GREATER THAN 25 MM (1 IN.). D, POSTWELD STRESS RELIEF IS REQUIRED FOR ALL REPETITIVELY LOADED OR SHOCK-LOADED STRUCTURES, FOR ALL RESTRAINED JOINTS, AND FOR ALL THICKNESSES OVER 50 MM (2 IN.); IT IS ALSO DESIRABLE FOR STATICALLY LOADED STRUCTURES. E, POSTWELD STRESS RELIEF IS RECOMMENDED FOR ALL APPLICATIONS; NO INTERMEDIATE COOLDOWN SHOULD BE PERMITTED FOR RESTRAINED STRUCTURES OR FOR BASE METAL HAVING A THICKNESS GREATER THAN 50 MM (2 IN.). F, SAME AS E, EXCEPT HAZARDS ARE GREATER

Bead Tempering. In a multipass weld, the underlying weld metal is subjected to repeated thermal cycles. This extends

the time for hydrogen to diffuse out of the weld area, and it tempers any martensite that may have previously formed. Each bead also provides the preheat for the subsequent weld pass. The final bead and the HAZ it creates will not receive this same benefit. If the base metal is crack-sensitive, the final bead should be placed so that the HAZ does not form in the base metal (Fig. 5).

FIG. 5 BEAD TEMPERING. THE LAST BEAD IS PLACED SO ITS HAZ IS FORMED IN THE WELD METAL RATHER THAN IN THE MORE HARDENABLE BASE METAL. SOURCE: REF 1

Use of Low-Hydrogen Processes and Consumables. The principal source of weld metal hydrogen is the welding

consumable(s). Careful selection of low-hydrogen consumables or low-hydrogen processes (such as gas-metal arc welding) can greatly reduce hydrogen pickup. This is especially important with medium- to high-carbon steels, in which the formation of a hard, brittle martensite is more likely. Many new specifications (for example, AWS A5.x filler metal specifications) now provide methods for consumable classification based on the diffusible hydrogen levels of their deposited weld metals. A standardized procedure is prescribed for measuring the diffusible hydrogen (for AWS specifications, this is AWS A4.3, while in many other parts of the world ISO 3690 is used). Designators are then used to indicate the maximum hydrogen level that an electrode will produce. The AWS uses a logarithmic scale; the International Institute of Welding method uses a linear system (Table 4).

TABLE 4 CLASSIFICATION OF CONSUMABLES RELATIVE TO THE DIFFUSIBLE HYDROGEN LEVELS OF THEIR DEPOSITED WELD METALS

SPECIFICATION HYDROGEN LEVELS AT INDICATED DESIGNATIONS, ML/100 G VERY LOW LOW HYDROGEN-CONTROLLED AWS 3 WT% SI NI3SI >2.5 WT% SI

NI3SI

>3 WT% SI

DOMINANT HARD PHASE NI3B NI3B AND CHROMIUM BORIDE (USUALLY CRB, ALTHOUGH CR2B AND CR3B2 MAY ALSO BE PRESENT) CRB AND CR5B3

WT%) ALL LEVELS

...

...

COMPLEX CARBIDES OF M23C6 AND M7C3 TYPES

FIG. 7 MICROSTRUCTURE OF ERNICR-C (TWO-LAYER PLASMA TRANSFERRED ARC DEPOSIT) BORIDECONTAINING NICKEL-BASE HARDFACING ALLOY. 425×. SOURCE: REF 2

The chief purpose of silicon in the material is to provide, in conjunction with boron, self-fluxing characteristics. However, as an important matrix element and as a potential promoter of intermetallic precipitates, it also has a powerful influence on the wear properties of the alloys. Boron content influences the level of silicon required for silicide (Ni3Si) formation. The higher the boron content, the lower is the silicon content required to form silicides. Because of the boride and carbide dispersions within their microstructures, the boride-containing nickel-base alloys exhibit excellent resistance to abrasion (ERNiCr-C, ERNiCr-B, and Alloy 40 in Fig. 4). Low-stress abrasion resistance generally increases with boron and carbon contents, hence the hard-phase volume fraction for these materials. Although their performance is not as good under self-mated sliding conditions as that of cobalt-base materials, the boridecontaining nickel-base alloys possess moderate resistance to galling. Of the nonferrous materials, the boride-containing nickel-base alloys are the least resistant to corrosion. This is attributed to the lack of chromium in the matrix that follows boride and carbide formation. Originally, the boride-containing nickel-base alloys were developed from brazing alloy compositions specifically for use with the spray-and-fuse process (for which their self-fluxing properties are ideal). During the fusing cycle of this process, it is believed that oxides within the sprayed coating combine with some of the silicon and boron to form a borosilicate slag (which floats to the surface of the deposit). The only other commercially significant spray-and-fuse alloys are some silicon-containing and boron-containing derivatives of the carbide-containing cobalt-base materials. Microstructurally, and in a property sense, these alloys bear little resemblance to their traditional counterparts because the presence of boron markedly changes the matrix chromium content and the types of hard phase that form during solidification. Additional information on the spray-and-fuse process can be found later in this article in the section on "Thermal Spraying." Carbide-Containing Alloys. The use of carbide-containing nickel-base alloys has been somewhat limited to date.

However, these alloys are gaining popularity as alternatives to cobalt-base hardfacing alloys. This is particularly applicable in the nuclear power industry, where it has been found that high-cobalt hardfacing alloys are sources of the highly radioactive 60Co isotope that deposits in pipes, fittings, and valve bodies (Ref 6). Wear and corrosion of hardfacing alloys lead to the release of 59Co into the primary coolant. The 59Co deposits on the reactor core, where it is transmuted to the radioactive 60Co isotope, which is then released back to the coolant. The most popular and widely used cobalt-free alloys are in the Ni-Cr-Mo-C and Ni-Cr-Fe-C systems. Table 6 lists several compositions of these alloys. Generally, carbide-containing nickel-base alloys have poor oxyacetylene weldability, which has made then unsuitable for hardfacing applications in the past. Carbide-containing nickel-base alloys of the Ni-Cr-Mo-Co-Fe-W-C system are more

readily weldable by oxyacetylene methods and are gaining popularity as low-cost alternatives to cobalt-base alloys. These alloys, depending on precise composition, contain M7C3- or M6C-type carbides similar to cobalt-base alloys. Laves Phase Alloys. Only one Laves phase-containing nickel-base alloy is commercially available: Ni-33Mo-16Cr-

3.5Si (T-700 in Table 6). This alloy, like most nickel-base alloys, is difficult to weld using the oxyacetylene process, but it can be readily welded using gas-tungsten arc welding or the PTA process. It can also be applied using the plasma or high-velocity oxyfuel thermal spray techniques. Although it has excellent metal-to-metal wear resistance and moderate abrasive wear resistance, it possesses poor impact resistance. High-Silicon Stainless Steels (Ref 2) Recently, several new high-silicon stainless steels have been introduced that may also be considered as alternate materials to cobalt-base alloys. These exhibit identical anti-galling characteristics and possess equal resistance to cavitation erosion. Some of these have been developed for general use and others for specific applications (for example, hydroelectric turbines and nuclear valves). Some are based on type 200 (manganese- and nitrogen-containing) austenitic stainless steel; others are based on type 300 stainless steel. One of these latter compositions is as follows (Ref 7):

ELEMENT CARBON CHROMIUM NICKEL SILICON MANGANESE IRON

COMPOSITION, WT% 0.12 21.0 8.0 5.0 6.5 BAL

These high-silicon hardfacing alloys differ from the traditional carbide-containing cobalt-base alloys in terms of their corrosion resistance (they are better in some aqueous media and worse in others), their mechanical properties (particularly at high temperatures), and their thermal stability. Copper-Base Alloys (Ref 2) Copper-base hardfacing alloys fall in the aluminum bronze category. Examples include ECu-AlB (Cu-9Al-4Fe-1Si) and ECuAl-D (Cu-13.5Al-4Fe). Silicon and other types of bronze are also available in welding consumable form for hardfacing. Aluminum, silicon, and iron are added to the bronzes to strengthen the solid solution, and, in excess of their solubility limits (about 8%, in the case of aluminum), to provide precipitation hardening. Wear tests have shown that the aluminum bronzes possess anti-galling characteristics (self-coupled) comparable to those of ER-CoCr-E; low-stress abrasion resistance of the aluminum bronzes, however, was very low. Components typically protected using the aluminum bronzes include gears, cams, bearings, and cold drawing/forming dies. These alloys are not recommended for elevated-temperature use because their mechanical properties considerably decrease at temperatures >200 °C (>390 °F). Hardfacing Alloy Selection (Ref 1) Hardfacing alloy selection is guided primarily by wear and cost considerations. However, other manufacturing and environmental factors must also be considered, such as base metal; deposition process; and impact, corrosion, oxidation, and thermal requirements. The factors affecting hardfacing process selection are discussed in the next section of this article. Usually, the hardfacing process dictates the hardfacing or filler-metal product form. Hardfacing alloys usually are available as bare rod, flux-coated rod, long-length solid wires, long-length tube wires (with and without flux), or powders. The most popular processes, and the forms most commonly associated with each process, are (Ref 2):

HARDFACING PROCESS OXYFUEL/OXYACETYLENE (OFW/OAW) SHIELDED METAL ARC (SMAW) GAS-TUNGSTEN ARC (GTAW) GAS-METAL ARC (GMAW) FLUX-CORED OPEN ARC SUBMERGED ARC (SAW) PLASMA TRANSFERRED ARC (PTA) LASER BEAM

CONSUMABLE FORM BARE CAST OR TUBULAR ROD COATED SOLID OR TUBULAR ELECTRODE) BARE CAST OR TUBULAR ROD TUBULAR OR SOLID WIRE TUBULAR WIRE (FLUX CORED) TUBULAR OR SOLID WIRE POWDER POWDER

ROD

(STICK

In general, the impact resistance of hardfacing alloys decreases as the carbide content increases. As a result, in a situation where a combination of impact and abrasion resistance is desired, a compromise between the two must be made. In applications where impact resistance is extremely important, austenitic manganese steels are used to build up worn parts. Frequently, wear is accompanied by aqueous corrosion from acids or alkalis, such as those encountered in the chemical processing or petroleum industries or in flue gas scrubbers. Few of the iron-base hardfacing alloys possess the necessary corrosion resistance in such aqueous media. As a result, nickel- or cobalt-base hardfacing alloys generally are recommended when corrosion resistance combined with wear resistance is required. For example, a knife used to cut tomatoes in a food processing plant can last many times longer than a tool steel knife if the edge is made of a cobalt-base alloy. Oxidation and hot corrosion resistance of iron-base alloys is also generally poor. Typically, boride-containing nickel-base alloys do not contain sufficient chromium in the matrix to resist oxidation. Hence, Laves phase-containing or carbidecontaining nickel- or cobalt-base alloys typically are recommended for applications in which wear resistance combined with oxidation or hot corrosion resistance is required. The ability of an alloy to retain strength at elevated temperatures is important for wear applications such as hot forging dies or valves for service at 870 °C (1600 °F), as well as service in coal gasification/liquefaction applications. Iron-base alloys with martensitic structures lose hardness at elevated temperatures. Generally, the high-temperature strength retention of a hardfacing alloy increases with its tungsten or molybdenum content. In applications requiring elevatedtemperature strength and wear resistance, cobalt-base alloys or Laves phase alloys are recommended. Wear-resistant materials typically are classified by type and alloy content, which largely determines their properties. These materials have structures consisting of hard phases (carbides, borides, or Laves phases) in a matrix hardened to a level that depends on the desired alloy properties. Table 9 provides an extensive list of hardfacing materials along with their specific advantages and applications. Such recommendations are also available from producers of hardfacing alloys. The following steps should be taken in selecting a hardfacing alloy: • • • • • •

ANALYSIS OF THE SERVICE CONDITIONS TO DETERMINE THE TYPE OF WEAR AND ENVIRONMENTAL RESISTANCE REQUIRED SELECTION OF SEVERAL HARDFACING ALLOY CANDIDATES ANALYSIS OF THE COMPATIBILITY OF THE HARDFACING ALLOYS WITH THE BASE METAL, TAKING INTO CONSIDERATION THERMAL STRESSES AND POSSIBLE CRACKING FIELD TESTING OF HARDFACED PARTS SELECTION OF AN OPTIMUM HARDFACING ALLOY, CONSIDERING COST AND WEAR LIFE SELECTION OF THE HARDFACING PROCESS FOR PRODUCTION OF WEAR COMPONENTS, CONSIDERING DEPOSITION RATES, THE AMOUNT OF DILUTION, DEPOSITION,

EFFICIENCY, AND OVERALL COST, INCLUDING THE COST OF CONSUMABLES AND PROCESSING

TABLE 9 SELECTION GUIDE FOR HARDFACING ALLOYS

TYPES OF MATERIALS LOW HARDNESS PEARLITIC STEELS

ADVANTAGES

APPLICATIONS

LOW COST, CRACK-RESISTANT

A BUILDUP MATERIAL TO RESTORE DIMENSION OR AS A LOW-COST BASE FOR SURFACING APPLICATIONS

MILD STEELS

AN EXCELLENT BASE FOR HARDFACING LOW-ALLOY SOME TOUGHNESS FOR STEELS BUILDUP OF WORN AREAS AUSTENITIC STEELS EXCELLENT FOR HEAVY IMPACT, TOUGH

FOR METAL-TO-METAL WEAR UNDER HEAVY IMPACT CONDITIONS LOW-CARBON NI- GOOD CORROSION RESISTANCE CORROSION-RESISTANT CR STAINLESS SURFACING OF LARGE TANKS, NUCLEAR VESSELS HIGH-CARBON OXIDATION AND HOT WEAR FURNACE PARTS, RED HEAT NI-CR STAINLESS RESISTANT FRICTIONAL WEAR 14% MN, CR-NI HIGH YIELD STRENGTH BUILDUP, CRACK REPAIR, JOINING AUSTENITIC MANGANESE ALLOYS OR MANGANESE TO MILD STEEL 14% MN, 1% MO FAIR ABRASION AND RAILWAY TRACK BUILDUP CORROSION RESISTANCE, WORK HARDENS 14% MN, 3% NI WORK HARDENS BUILDUP AND JOINING OF MANGANESE PARTS SEMIAUSTENITIC LOW-COST CRACK-RESISTANT HARDFACING APPLICATIONS STEELS MATERIALS COPPER ALLOYS RESISTANCE TO SEIZING FOR BUILDUP OF BEARING UNDER FRICTIONAL WEAR SURFACES CONDITIONS MEDIUM HARDNESS NICKEL-BASE GOOD CORROSION RESISTANCE FOR APPLICATIONS WHERE HOT ALLOYS AND EXCELLENT HOT HARDNESS IS IMPORTANT HARDNESS NICKELRESISTANCE TO OXIDATION FOR CORROSIVE WEAR CHROMIUM CONDITIONS NICKELRESISTANCE TO EXHAUST GAS FOR HARDFACING TRUCK, BUS, CHROMIUMCORROSION AND AIRCRAFT ENGINE VALVES MOLYBDENUM NICKELCORROSION-RESISTANT FOR SURFACING APPLICATIONS CHROMIUMPROPERTIES MOLYBDENUMTUNGSTEN NICKELCORROSION-RESISTANT AS PUMPING SERVICE APPLICATIONS, CHROMIUM-BORON WELL AS ABRASIVEESPECIALLY IN THE OIL FIELDS

MARTENSITIC ALLOY IRONS

RESISTANT PROPERTIES EXCELLENT ABRASION RESISTANCE

CHROMIUMMOLYBDENUM

ABRASION RESISTANCE WHILE WITHSTANDING LIGHT IMPACT

CHROMIUMTUNGSTEN AUSTENITIC ALLOY IRONS

HIGH COMPRESSIVE STRENGTHS BETTER CRACK-RESISTANT PROPERTIES THAN THE MARTENSITIC ALLOY IRONS GOOD CRACK-RESISTANT QUALITIES CRACK-RESISTANT, LIGHT IMPACT ABRASION RESISTANCE WITH MEDIUM IMPACT

CHROMIUMMOLYBDENUM NICKELCHROMIUM MARTENSITIC STEELS

LOW CARBON (UP TO 0.30% C) MEDIUM CARBON (0.30-0.65% C) HIGH CARBON (0.65-1.7% C) HIGH HARDNESS HIGH-CHROMIUM IRONS

FOR OVERCOMING ABRASIVE WEAR WHERE LIGHT OR NO IMPACT IS INVOLVED FOR THE BUILDUP OF MACHINE PARTS SUBJECT TO REPETITIVE METAL-TO-METAL WEAR WITH OR WITHOUT LIGHT IMPACT

WHERE EROSIVE WEAR WITH OR WITHOUT LIGHT IMPACT IS PRESENT PUMP AND TURBINE BUILDUP FOR EROSIVE WEAR WITH LIGHT IMPACT FOR A VARIETY OF ABRASIVE WEAR CONDITIONS WHERE THERE MAY ALSO BE INVOLVEMENT WITH MEDIUM IMPACT

LEAST EXPENSIVE, TOUGH GOOD RESISTANCE TO MEDIUM IMPACT FAIR ABRASION RESISTANCE

EXCELLENT EROSIONRESISTANT PROPERTIES

AUSTENITIC

GOOD ABRASION-RESISTANT QUALITIES

MARTENSITIC

ABILITY TO BE REHARDENED AFTER ANNEALING

TUNGSTENMOLYBDENUM ALLOYS

HIGH RED HARDNESS

CHROMIUMTUNGSTEN-COBALT ALLOYS

HIGH RED HARDNESS, HIGH HOT STRENGTH AND CREEPRESISTANT QUALITIES

LOW CARBON (UP TO 1.0% C) MEDIUM

GOOD OXIDATION RESISTANCE, TOUGH RESISTANT TO ABRASION AND

FOR GENERAL USE WHERE HOT GASES OR MATERIALS ARE INVOLVED FOR FARM AND EARTH-MOVING EQUIPMENT WORKING IN SANDY SOILS AS A HARDFACING IN STEEL MILLS AND REFINERIES WHERE HOT EROSION (595 °C, OR 1100 °F) IS A WEAR PROBLEM FOR HARDFACING OF COKE OVEN PARTS AS WELL AS OTHER HOT (425 TO 650 °C, OR 800 TO 1200 °F) STEEL MILL APPLICATIONS FOR MOST HARDFACING APPLICATIONS, BUT BEST SUITED TO THOSE APPLICATIONS WHERE HOT WEAR AND ABRASION (ABOVE 650 °C, OR 1200 °F) ARE INVOLVED (JET ENGINE TURBINE BLADES, GAS ENGINE EXHAUST VALVE)

CARBON (1.4% C) HIGH CARBON (2.5% C) TUNGSTEN CARBIDES

OXIDATION GOOD ABRASION RESISTANCE, BRITTLE THE ULTIMATE IN ABRASIONRESISTANT QUALITIES

FINE-GRAIN TUBULAR RODS

A RELATIVELY SMOOTH, EXTREMELY WEAR-RESISTANT SURFACE, GAS WELD DEVELOPS A ROUGH CUTTING SURFACE AS THE DEPOSITED BINDER WEARS AWAY DEVELOPS AN EXTREMELY ROUGH ABRASIVE CUTTING SURFACE AS GRANULES OR INSERTS BECOME MORE EXPOSED. HIGHLY ABRASIVE, BUT BRITTLE UNDER IMPACT CONDITIONS

COARSE GRANULAR TUBE ROD GRANULES OR INSERTS

A VARIETY OF MATERIALS USED FOR HARDFACING TO MEET AN EXTREMELY WIDE RANGE OF SEVERE ABRASIVE CONDITIONS, ESPECIALLY OIL WELL DRILL BITS, TOOL JOINTS, ROCK DRILL BITS

Source: Ref 8

Hardfacing Process Selection (Ref 1) Hardfacing process selection may be as important as hardfacing alloy selection, depending on the engineering application. Service performance requirements not only dictate hardfacing alloy selection, but also have a strong influence on hardfacing process selection. Other technical factors involved in hardfacing process selection include (but are not limited to) hardfacing property and quality requirements, physical characteristics of the workpiece, metallurgical properties of the base metal, form and composition of the hardfacing alloy, and welder skill. Ultimately, economic considerations predominate, and cost is the determining factor in the final process selection. The term "hardfacing" is sometimes used to describe true welding processes, in which both the filler material and a small amount of the base metal are molten and are bonded metallurgically. Such processes, which include oxyfuel, arc, and laser welding methods, are preferred for applications that require dense and relatively thick coatings with excellent bonding between the hardfacing material and the base metal. Thermal spray processes are used for applications that require a thin, hard coating applied with minimum work distortion and where a purely mechanical bond between the hardfacing material and the base metal is acceptable. Property and Quality Requirements. Welded hardfacing deposits are, in effect, mini-castings characterized by variable composition (segregation) and solidification kinetics that influence deposit microstructure. It is not surprising, therefore, that the properties and quality of welded hardfacing deposits should depend on welding process and technique, as well as on alloy selection.

Composition variations also derive from base-metal dilution during welding, although carbon pickup in oxyfuel gas welding (OFW) processes and solute volatilization during arc welding processes are additional factors. Dilution is the inter-alloying of the hardfacing alloy and the base metal and is usually expressed as the percentage of base metal in the hardfacing deposit. A dilution of 10% means that the deposit contains 10% base metal and 90% hardfacing alloy (see the section on "Weld Cladding" in this article for more detailed information). The wear resistance and other desirable properties of the hardfacing alloy are generally thought to degrade as dilution increases. The maximum amount of allowable dilution depends on specific service requirements. Welding process and technique should, however, be selected so as to control dilution to less than 20% for the hardest layer. More dilution is acceptable for buildup layers. Dilution is nil in thermal spraying processes and tends to be acceptably low in conventional OFW processes. Dilution is generally more of a problem in arc weld hardfacing processes, ranging from approximately 5% in plasma arc welding

(PAW) hardfacing processes to as much as 50% in submerged arc welding (SAW). Laser hardfacing dilution is also low, ranging generally from 1 to 10%. High-dilution welding processes can be tolerated in applications requiring relatively thick hardfacings that can be applied in multiple layers. The effects of dilution on the microstructure of a high-chromium, white iron hardfacing alloy deposited in multiple layers on low-carbon steel substrate using shielded metal arc welding (SMAW) are illustrated in Fig. 8. The first layer (Fig. 8a) microstructure in this particular case was hypoeutectic with interdendritic eutectic and exhibited a hardness of 55 HRC. The second layer (Fig. 8b) consisted of a fine hypereutectic microstructure with a large percentage of eutectic and exhibited a hardness of 57 HRC. The third (Fig. 8c) and fifth (Fig. 8d) layers were microstructurally equivalent, consisting of primary M7C3 carbides in eutectic matrices. The hardness in the third and subsequent layers ranged from 60 to 61 HRC.

FIG. 8 EFFECT OF DILUTION ON THE MICROSTRUCTURE OF A FE-28CR-4MO-0.4MN-4.6C HARDFACING ALLOY. (A) FIRST LAYER OVER MILD STEEL SHOWING A HYPOEUTECTIC STRUCTURE OF METAL DENDRITES AND INTERDENDRITIC EUTECTIC, 55 HRC. (B) SECOND LAYER SHOWING A FINE HYPOEUTECTIC STRUCTURE WITH A LARGE PERCENTAGE OF EUTECTIC MATRIX, 57 HRC. (C) THIRD LAYER SHOWING A COARSE HYPEREUTECTIC STRUCTURE OF PRIMARY M7C3 CARBIDES IN A EUTECTIC MATRIX, 61 HRC. (D) FIFTH LAYER SHOWING A STRUCTURE SIMILAR TO BOTH THE THIRD AND FOURTH PASSES, 60 HRC. SOURCE: REF 1

The microstructures and mechanical properties of hardfacing deposits, like castings, vary depending on solidification kinetics as well as on dilution. Solidification generally is rapid in thermal spray processes, with nearly amorphous structures being achieved with some thermal spray processes. Solidification kinetics tend to be somewhat slower in conventional weld hardfacing processes. As a rule, hardfacing deposits produced by OFW processes tend to solidify more slowly than hardfacing deposits produced by arc welding processes. These differences in solidification rate produce widely different microstructures and widely different properties regardless of dilution. The effects of welding process on the microstructure of a cobalt-base (ERCoCr-A) hardfacing alloy deposited on a low-carbon steel substrate are illustrated in Fig. 9.

FIG. 9 EFFECT OF WELDING PROCESS ON THE MICROSTRUCTURE OF AN ERCOCR-A HARDFACING ALLOY. (A) OXYFUEL GAS WELDING. (B) PLASMA TRANSFERRED ARC WELDING. (C) GAS-TUNGSTEN ARC WELDING. (D) FLUX-CORED OPEN ARC WELDING. (E) SUBMERGED ARC WELDING. (F) SHIELDED METAL ARC WELDING. SOURCE: REF 1

Physical Characteristics of the Workpiece. The size, shape, and weight of the workpiece have a strong influence on hardfacing process selection. Large, heavy workpieces that are difficult to transport are most conveniently hardfaced using manual or semiautomatic hardfacing processes in which the hardfacing equipment can be moved to the workpiece. Gas shielded metal arc and open arc welding processes, for which portable equipment is readily available, are well suited to field applications involving inaccessible surfaces on large, difficult-to-transport workpieces. Thermal spraying processes and equipment also are available for in situ field hardfacing applications.

Hardfacing process selection is more involved in applications where the workpiece is sufficiently small so that it can be readily transported to the welding equipment. Thermal spraying processes and OFW, SAW, gas-tungsten arc welding (GTAW), PAW, and gas-tungsten arc welding (GMAW) are all well suited to in-plant hardfacing of small workpieces. Preference is usually given to thermal spraying and/or OFW or GTAW for small workpieces requiring thin, accurately placed hardfacing deposits. Large quantities of small workpieces can usually be hardfaced most economically using dedicated, specially designed, fully automatic hardfacing equipment and processes.

Metallurgical Characteristics of the Base Metal. As described below, base-metal surface preparation is important in

thermal spray hardfacing processes. However, base-metal composition, melting temperature range, and thermal expansion and contraction characteristics have a significant effect on welding process selection. Steels are generally suitable base metals for hardfacing. Low-alloy and medium-carbon steels with carbon contents up to 0.4% can be easily hardfaced by all welding processes with generally excellent results. Higher-carbon steels are similarly readily amenable to hardfacing, but preheating is necessary to minimize deleterious martensitic reactions in the heataffected zone (HAZ). Austenitic stainless steels, with the exception of the free-machining grades, and most nickel-base alloys can also be easily hardfaced by all welding processes with excellent results. Martensitic stainless steels, tool steels, die steels, and cast irons are also amenable to hardfacing, but greater attention to preheat, interpass temperature, and postweld heat treatment is required. Age-hardenable base metals present special hardfacing problems and generally require solution and/or overaging heat treatments prior to hardfacing, as well as careful attention to preheat, interpass temperature, and postheat. Hardfacing welding processes require that the base metal melting temperature range be higher than or at least equal to that of the hardfacing alloy. Thermal spraying processes, on the other hand, do not require a base metal melting temperature advantage. Base metal thermal expansion and contraction characteristics affect solidification strain. Workpieces frequently have to be uniformly heated to relatively high temperatures during welding to enable the deposition of crack-free hardfacing alloys. Welding process selection is important in that it directly affects heat input and preheat requirements. Thermal expansion and contraction differences between the hardfacing alloy and base metal are important in applications involving thermal cyclic service conditions. Large differences can result in shear failures in thermal spray deposits or thermal fatigue failures in welded deposits. Buffer layers are often deposited between the base metal and hardfacing alloy to counteract large differences in thermal expansion and contraction characteristics. Frequently, dilution, which is generally regarded to be detrimental, can be used to good advantage to form in situ buffer layers and to minimize the risk of thermal fatigue failure. Hardfacing process selection is an important consideration in this regard. Hardfacing Product Forms. Hardfacing process selection is often limited by product form. Hardfacing alloys normally

are available as bare cast rod, covered electrodes, wire, and powders. However, not all alloys are available in all product forms. Virtually all hardfacing alloys can be produced in powder form, and most hardfacing alloys can be produced in rod form. Rods generally are available in bare cast form or in tubular form. Bare cast rod diameters are limited to a large extent by alloy characteristics and casting processes to a range of 3.2 to 8.0 mm (

1 5 to in.). Bare cast rod lengths usually range 8 16

from 250 to 700 mm (10 to 28 in.), but some hardfacing alloys can be cast in extra long, continuous lengths with diameters as small as 3.2 mm (

1 in.). Tubular rods consist of highly alloyed powder cores encased in low-alloy sheaths. 8

Tubular rods can be produced in discrete lengths or in spooled continuous coils. Discrete-length tube rods generally are available in the same diameter range as bare cast rod. Long-length tube rods, or tube wire as they are properly named, normally are available in diameters ranging from 4.0 to 1.6 mm (

5 1 to in.). Some of the tougher, more ductile 32 16

hardfacing alloys can be produced as solid drawn wire. Some of the more brittle hardfacing alloys can be produced, at a premium cost, as solid wire using extrusion techniques. Welder Skill. It is essential to consider hardfacing quality requirements in relation to welder/operator skill, as well as in

relation to the hardfacing process. As a rule, hardfacing by manual welding processes such as OFW or GTAW requires high welder proficiency, whereas hardfacing by automatic welding processes, such as SAW, requires a minimum of operator skill, assuming machine control settings have previously been established. Thermal spraying and related powder welding processes generally require intermediate welder/operator skills. Manual GTAW can be used to obtain high-quality deposits on relatively small areas, such as the interlock regions of turbine blades. In this type of application, thin layers can be deposited with dilution as low as 10%, but relatively high welder skill and close control of the welding operation are necessary. In contrast, earthmoving and mining equipment can be hardfaced adequately in the field by relatively unskilled welders. Selection of process is usually based on maximum deposition rate: high dilution rates seldom significantly affect the suitability of the coating for service.

Cost is ultimately the determining factor in hardfacing process selection. Hardfacing costs basically consist of labor,

materials, and in some instances, new equipment cost. Transportation cost also must be considered. Labor costs depend primarily on the level of welder skill required and process deposition rates, but surface preparation and finishing costs also must be considered. Deposition rates are normally lowest with manual hardfacing processes. Considerably higher deposition rates are possible with automatic arc welding hardfacing processes. Deposition rates as high as 27 kg/h (60 lb/h) are possible with multiwire or powder and wire submerged arc hardfacing processes, but dilution is also high. An ideal welding process for hardfacing would be one with the ability to limit dilution while achieving a high deposition rate. Typical dilution percentages; deposition rates; and minimum deposit thicknesses for different welding processes and various forms, compositions, and modes of application of hardfacing alloys are given in Table 10.

TABLE 10 CHARACTERISTICS OF WELDING PROCESSES USED IN HARDFACING WELDING PROCESS

OAW

MODE OF APPLICATION

MANUAL

MANUAL AUTOMATIC

SMAW

MANUAL

OPEN ARC

SEMIAUTOMATIC

AUTOMATIC

GTAW

MANUAL

AUTOMATIC SAW

AUTOMATIC, SINGLE WIRE AUTOMATIC, MULTIWIRE

FORM OF HARDFACING ALLOY

WELDMETAL DILUTION, %

kg/h

lb/h

MINIMUM THICKNESS(A) mm in.

0.52

1-4

0.8

MATH OMITTED

100

0.52 0.57

1-4

0.8

1 32 1 32

85-95

10-20

0.55

1-12 3.2

1 8

65

15-40

2-11

5-25 3.2

1 8

80-85

15-40

2-11

5-25 3.2

1 8

80-85

10-20

0.53

1-6

2.4

3 32

98-100

10-20

0.55 5-11

1-10 2.4

98-100

1025

3.2

3 32 1 8

1127

2560

4.8

3 16

95

BARE CAST 1-10 ROD, TUBULAR ROD POWDER 1-10 EXTRA-LONG BARE CAST ROD, TUBULAR WIRE FLUXCOVERED CAST ROD, FLUXCOVERED TUBULAR ROD ALLOYCORED TUBULAR WIRE ALLOYCORED TUBULAR WIRE BARE CAST ROD, TUBULAR ROD VARIOUS FORMS(B) BARE TUBULAR WIRE BARE TUBULAR

DEPOSITION

1-10

30-60

15-25

1-15 0.8

DEPOSIT EFFICIENCY, %

100

95

WIRE BARE TUBULAR WIRE POWDER(C)

AUTOMATIC, SERIES ARC PAW

AUTOMATIC MANUAL

AUTOMATIC GMAW

SEMIAUTOMATIC

AUTOMATIC

LASER

AUTOMATIC

BARE CAST ROD, TUBULAR ROD VARIOUS FORMS(B) ALLOYCORED TUBULAR WIRE ALLOYCORED TUBULAR WIRE POWDER

10-25

1116

2535

4.8

3 16

95

5-15

0.57 0.54

1-15 0.8

85-95

1-8

2.4

1 32 3 32

0.54 0.95

1-8

2.4

3 32 1 16

98-100

10-40

0.95

2-12 1.6

1 16

90-95

1-10

(D)

(D)

5-15

5-15 10-40

2-12 1.6

0.13 0.005

98-100

90-95

85-95

(A) RECOMMENDED MINIMUM THICKNESS OF DEPOSIT. (B) BARE TUBULAR WIRE; EXTRA-LONG (2.4 M, OR 8 FT) BARE CAST ROD; TUNGSTEN CARBIDE POWDER WITH CAST ROD OR BARE TUBULAR WIRE. (C) WITH OR WITHOUT TUNGSTEN CARBIDE GRANULES. (D) VARIES WIDELY DEPENDING ON POWDER FEED RATE AND LASER INPUT POWER Hardfacing material costs depend on the prevailing price of raw materials and product form. Raw material cost is predominant in tungsten-, cobalt-, and nickel-base hardfacing alloys, whereas product form is the predominant cost in iron-base hardfacing alloys. As a rule, tubular rod and wire are the least expensive hardfacing product forms. In contrast, solid wire is the most expensive hardfacing product form. Cast rod and powder are generally intermediate in cost. Powder, depending on specific powder size requirements, tends to be less expensive than cast rod. However, deposit efficiency, measured as the percentage of hardfacing consumable retained on the workpiece surface, is generally slightly lower for hardfacing processes that utilize powder consumables than for hardfacing processes that utilize rod consumables. Thermal spraying processes are particularly poor in this regard, with typical deposit efficiencies as low as 70%. Equipment costs vary from simple arc welding equipment to fully automated laser beam hardfacing systems. The more advanced automatic hardfacing equipment often costs 100 to 1000 times as much as simple arc welding equipment, depending on the level of instrumentation and computer control. It is relatively easy to estimate or predetermine the cost of hardfacing. It is more difficult to estimate the value of hardfacing. For maintenance applications, it has been suggested that hardfacing is preferred over total replacement when the cost advantage (CA) is positive. The cost advantage can be calculated by: CA =

CN CR − PN PR

(EQ 1)

where CN is the cost of a new component; CR is the cost of hardfacing plus downtime costs; PN is the work output during the service life of a new component; and PR is the work output during the service life of a hardfaced component. Most of the data needed for the above calculation usually can be obtained; PR is perhaps the most difficult value to assess, but reasonable estimates can be made based on relative laboratory wear test results. If the workpiece is hardfaced with an alloy of essentially the same chemical composition as the wom material, it is reasonable to assume that PR = PN. On the

other hand, if the hardfacing material is more wear-resistant than the worn material, then PR > PN. Therefore, from Eq 1, CA is likely to be positive, which is the primary objective of hardfacing. Hardfacing by Oxyfuel Gas Welding Oxyfuel gas welding (OFW) is a process in which the heat from a gas flame is used to melt hardfacing materials onto workpiece surfaces. During hardfacing, the workpiece is only superficially melted. The gas flame is created by the combustion of a hydrocarbon gas and oxygen. Acetylene is the most widely used fuel for OFW. In fact, OFW is commonly referred to as oxyacetylene welding (OAW), even though other fuel gases, such as methylacetylene propadiene (MAPP), occasionally are used in certain OFW processes. Hardfacing by OFW is a versatile process, readily adaptable to all hardfacing product forms, including powder. As described below, OFW hardfacing processes that use fillers in powder form are commonly referred to as manual powder torch welding or powder welding. Oxyfuel gas welding equipment is simple, and capital equipment costs are low. Furthermore, OFW equipment is mobile and can be used for fusion welding and cutting operations, as well as for hardfacing. More importantly, OFW processes generally produce hardfacing deposits with optimum microstructures for wear-resistant applications. Hardfacing by OFW, however, is not without limitations. Hardfacing by manual OAW requires a high degree of welder skill to obtain smooth, high-quality deposits, as the welding rod and the torch flame have to be separately manipulated. The speed of welding and low deposition rates often preclude the use of OAW for hardfacing applications on large components that require large quantities of hardfacing alloy, unless automation of the operation is warranted by configuration of the surface and a large number of identical parts. Oxyacetylene welding is best suited to hardfacing applications in which the surface area to be hardfaced is minimal. Thus, for hardfacing steam valves, automotive and diesel engine valves, cutter blades for wood and plastic, chain saw bars, and plowshares and other agricultural implements, it is often the most satisfactory process as it can be easily employed in field applications. Most steel workpieces can be hardfaced by OFW processes. High-manganese and high-sulfur steels are exceptions. Highspeed steel workpieces are difficult to hardface and should be fully annealed before hardfacing. Cast iron workpieces also can be hardfaced by OFW processes, but cast iron substrates generally require the use of fluxes and special welding precautions. Dilution is generally low (1 to 10%) in OFW hardfacing deposits, as only the surface of the workpiece in the immediate area being hardfaced is brought to the melting temperature. The hardfacing rod is melted and spread over the surface of the workpiece. The melted rod does not mix or alloy with the base metal, but only with the melted surface. Superficial workpiece melting of steel workpieces produces a wet appearance, called "sweating." Operating Procedures. Workpiece surfaces must be thoroughly cleaned before hardfacing. If impurities are present on

surfaces of the workpiece, or if an excessive amount of oxide is produced in heating (usually, a neutral or reducing flame is used), it may be necessary to dislodge the solid contaminants by prodding or rubbing with a hardfacing rod tip. Oxide trapped beneath the hardfacing overlay can react with carbon in the hardfacing alloy to produce a gas and thus cause porosity. Hardfacing alloys usually contain a deoxidizer that controls a moderate amount of oxidation. It is often necessary to preheat workpieces prior to hardfacing, regardless of the welding process, to prevent cracks in the hardfacing deposit. Preheating not only reduces tensile strain associated with thermal contraction differences between the hardfacing alloy and the workpiece, but in the case of OFW, also permits the use of a softer flame and results in less dilution. Small low-carbon steel workpieces can be preheated locally with a torch, but large workpieces are best preheated with gas burners in a brick muffle furnace, allowing sufficient soak time to ensure even heating. With practice, one-layer hardfacing deposits up to 1.6 mm (

1 in.) thick can be deposited, particularly if the workpiece is 16

inclined so that the deposition is slightly uphill. Alternatively, thick deposits can be applied in several layers. Hardfacing deposits by OFW can be remelted as necessary to improve deposit quality or surface appearance. Common Flaws. Hardfacing deposits applied by OFW generally are characterized by low dilution and high deposit

quality. Discontinuities are possible, however, depending on the skill of the welder and the quality of the filler material.

Porosity is probably the most common flaw observed in hardfacing deposits applied by OFW. Porosity in hardfacing deposits, like porosity in castings, can result from inadequate feeding or gas evolution during solidification. Porosity resulting from inadequate feeding is sometimes referred to as microshrinkage. Microshrinkage is most common at the termination of the hardfacing deposit and usually results from the inability of the hardfacing deposit to accommodate solidification shrinkage after the flame is withdrawn. This type of flaw can be minimized or prevented by withdrawing the flame very slowly and allowing solidification to progressively occur from the bond line to the surface. Porosity can also occur by gas evolution during solidification if the total gas content in the molten deposit is greater than the solid solubility limit. Gas-induced porosity generally derives from slag or non-metallic inclusions trapped in or on the surface of the base metal. However, the initial gas content of the filler materials and possible gas pickup during OFW are also contributing factors. Manual Powder Torch Welding (Ref 9) Also termed "powder welding," manual powder torch welding is a modified OFW process in which powder application and fusion occur in a single operation using a special oxyacetylene torch. Powder is fed from a small hopper mounted on a gas welding torch into the fuel gas supply and conveyed through the flame to the workpiece surface. By following general oxyacetylene surfacing practice, smooth, thin, dense deposits are obtained. Deposit thickness is controlled by powder flow rate and movement of the torch. A dilution of 1 to 5% in the deposit occurs from this process. Deposits ranging from 0.75 to 3 mm (0.030 to 0.125 in.) in thickness can be applied at a rate of 1.3 mm (0.050 in.) per pass. Sweating of the base metal during surfacing generates a fusion-type bond with the hardfacing alloy. Ease of application and material recovery are enhanced because of close control of the powder flow into the welding flame. Oxyacetylene torch flame adjustment is essential to produce good powder deposition. The amount of excess acetylene required for hardfacing varies depending on alloy composition. The amount of excess acetylene is conveniently measured in terms of the flame geometry as the ratio between the length of the oxyacetylene feather and the length of the inner cone. The length of the inner cone, indicated by the value X, is the distance from the torch tip to the extreme end of the inner cone. The length of the oxyacetylene feather is measured from the torch tip to the extreme end of the intermediate zone, generally expressed in multiples of the length of the inner cone. A neutral flame is thus classified as 1×. Cobalt-base powders are usually applied using 2× to 3× flame. The oxyacetylene feather is two to three times the length of the inner cone. Tungsten carbide-containing composite powders should be applied using a 1.5× flame. Nickel-base powders are welded with a neutral or slightly oxidizing flame. Because flame characteristics change somewhat when powder is being sprayed, the powder feed handle should be depressed and final flame adjustment made while the powder is being sprayed. Alloys commonly welded with this process contain boron and silicon, which combine with any oxide skin to form an easily removed slag and to provide a clean weld pool. The manual torch process is well suited for small repairs or for hardfacing of small areas requiring a wide range of hardnesses (about 20 to 60 HRC). Typical applications of manual torch welding include glass dies and molds, brick molding equipment, components for the cement industry, cast iron parts for foundries, feed screws, hammer mill hammers, and pump parts. Hardfacing by Arc Welding (Ref 1) In hardfacing by arc welding, the heat from an electric arc is used to melt the hardfacing material onto the workpiece surface. The electric arc is developed by impressing a voltage between an electrode and the workpiece. The voltage required to sustain the arc varies with the distance between the electrode and workpiece and with the arc welding process. The filler material is provided by the electrode used to form the arc or can be externally introduced into the arc. Arc welding processes in which the filler derives directly from the electrode are sometimes referred to as consumable electrode processes. Arc welding processes in which the filler is externally introduced into the arc are sometimes referred to as nonconsumable electrode processes. As a rule, nonconsumable electrode processes are accomplished at lower power requirements than are consumable electrode processes, resulting in less dilution and lower deposition rates. Consumable and nonconsumable electrode hardfacing processes both require that the filler materials be protected from oxidation as they are melted and joined to the workpiece surfaces.

Filler materials in consumable electrode hardfacing processes are sometimes protected by fluxes as in SMAW, SAW, and flux-cored open arc welding, or by inert gases, as in GMAW. The selection of a particular consumable or nonconsumable electrode hardfacing process depends on consideration of the factors discussed in the section on "Hardfacing Process Selection" in this article. Shielded Metal Arc Welding. The electrodes used in the SMAW consumable electrode hardfacing process generally consist of discrete lengths of filler material rod covered with specially formulated coatings. The filler material rod

generally ranges in diameter from 3.2 to 8.0 mm (

1 5 to in.) and may be cast or of tubular construction. The electrode 8 16

coating protects the filler material during welding by chemical reaction as the electrode melts. The coating can be fortified to increase alloy content during melting, as well as to simply provide molten metal protection. Coating ingredients also determine or affect the shape of the weld, amount of penetration, cooling rate of the weld, stabilization of the arc, and refinement of the hardfacing deposit. The selection of the proper power supply for SMAW consumable electrode hardfacing is not a serious problem. Motorgenerators, direct current rectifiers, and alternating current transformers can be used. Direct current electrode positive (DCEP) normally is recommended for hardfacing. Direct current electrode negative (DCEN) can be used to increase deposition rates slightly, but there is a tendency for the weld beads to be high and narrow. Although direct current is recommended, alternating current has been used in many applications with reasonably good success. Arc stability and burnoff characteristics are not quite as good by this method. Speed of travel, position of the electrode in relation to the weld pool, arc voltage, and amperage are modified to minimize 1 in.) diam 8 1 hardfacing electrodes range from 60 to 130 A, depending on the alloy. Welding currents for larger 6.4 mm ( in.) diam 4

penetration and the consequent dilution of the deposit. Typical welding currents for small 3.2 mm (

hardfacing electrodes range from 150 to 330 A. Depending on the surface area to be hardfaced, the bead-deposit pattern may be of the stringer or weave type, with or without staggering of beads, as dictated by service requirements. Selection of the bead pattern and sequence also depends on the accessibility of the surface to be hardfaced and the distortion limits of the workpiece. Deposition rates vary considerably with welding conditions. Normally from 1.4 to 1.8 kg (3 to 4 lb) of metal per hour of arc time can be applied. The recovery of alloy in the deposit generally ranges from 65 to 75% of the electrode weight, excluding stub loss. The average stub loss is approximately 75 mm (3 in.), which corresponds to about 20% of a 350 mm (14 in.) long electrode. There are many advantages to the SMAW consumable electrode process. The intense heat of the arc permits the hardfacing of large parts without preheat. Small areas of parts machined or fabricated to close tolerances can be hardfaced without distorting the entire part. Difficult-to-reach areas are best accessed with coated electrodes. The overall process is relatively fast, as preheat and high interpass temperatures do not have to be maintained. Another attractive feature of this process is the portability of the equipment. Gasoline engine-driven generators can be readily taken into the field. There are also some disadvantages to the SMAW consumable electrode process that must be considered. Penetration tends to be high, and dilution rates of 20% or more are not uncommon. This results in lower hardness and less resistance to abrasion. Two or three layers of weld metal normally are required to obtain maximum wear properties. Coated electrodes also leave a slag covering on the surface of the weld deposit which must be removed prior to each succeeding pass. Because coated electrodes usually are deposited with no more than a small amount of preheat, cracking or cross checking may occur, especially with the higher alloy electrodes that produce hard overlays. The hardfacing material is at its maximum temperature, and hence its maximum thermal expansion, as it is being deposited. The workpiece, by contrast, may be near ambient temperature during the welding process and may be considerably less expanded than the hardfacing material. Severe strains develop in the hardfacing deposit as it cools after welding. These strains are accommodated by plastic deformation, in the case of a ductile hardfacing material or by cracking and cross checking, in the case of more brittle hardfacing materials. Cracking or cross checking does not affect the wear resistance of the overlay in many applications.

Submerged arc welding is a consumable electrode hardfacing process in which the area between the electrode and

workpiece is shielded by a blanket of granular, fusible flux material. Electrodes in the SAW hardfacing process consist of continuous lengths, or coils, of filler materials--generally in tubular wire form. The fluxes used for hardfacing by SAW are mineral compositions specially formulated to protect the molten filler metal without appreciable gas evolution. Submerged arc fluxes are generally finely divided, free-flowing granules in the solid state, which are placed by gravity along the surface to be hardfaced in advance of the electrode. The flux becomes molten and highly conductive under the arc, creating localized conditions conducive to high deposition rates. Hardfacing filler materials can be deposited over a wide range of welding currents, voltages, and travel speeds, each of which can be controlled independently. Generally, however, the current, voltage, and workpiece travel speed are adjusted for hardfacing applications to produce a bead approximately 15 mm (

5 in.) wide with a slightly convex surface and sides 8

tapering into the base metal. Welding current usually ranges from 300 to 500 A for most SAW hardfacing applications. However, it is not uncommon to increase current to as much as 900 A to increase deposition rate. A number of features make SAW particularly attractive for hardfacing. Because it is an automatic process, SAW requires little skill on the part of the operator for efficient operation. High deposition rates are possible, usually ranging as high as 4.5 to 9.0 kg/h (10 to 20 lb/h), depending on the job. The process is capable of making porosity-free deposits that are smooth and even on the surface and that either can be used aswelded or require only a minimum of grinding to clean up. The filler metal is protected from the atmosphere by flux, which also completely surrounds the arc, requiring no shielding for the operator. The major disadvantages of hardfacing by SAW are that: (1) it is generally limited to hardfacing simple cylindrical or flat workpieces; and (2) it is not particularly adaptable to the surfacing of small parts. Hardfacing in other than the flat position requires special flux retainers. Preheating the workpiece in excess of 315 °C (600 °F) usually makes slag removal extremely difficult unless the slag is designed for easy hot slag removal. Occasionally, small portable units are used, but normally SAW units are fairly large and are mounted on stationary tracks, thus limiting their use for field operations. It is imperative that the flux be kept dry prior to use; therefore, storage can be a problem. Dilution is relatively high. It is not uncommon to find interalloying with the base metal, reaching or exceeding levels of 30%. Consequently, two or three layers of weld metal may be required for maximum wear resistance. Dilution can be reduced by long stickout techniques and by heavy overlap of individual weld beads. Flux-cored open arc welding is a semiautomatic or automatic process in which the arc between the electrode and

workpiece is shielded by a self-generated or self-contained gas or flux. Flux-cored open arc electrodes are continuous lengths, or coils, of filler material in tubular wire form. Open arc tubular filler materials are generally of low-carbon steel sheath construction with alloying elements, deoxidants, arc stabilizers, and a shielding gas source and/or flux contained in the core as powder or granules. High-alloy filler materials are generally shielded by a self-generated CO2 gas deriving from the reaction between carbonaceous materials contained in the core and the atmosphere, or by vaporization of fluorides in the corn. Arc shielding is often supplemented in low-alloy filler materials with flux and/or aluminum or magnesium additions, which serve as deoxidizers. Low equipment cost is one of the primary advantages of the flux-cored open arc hardfacing process. Neither fluxhandling nor gas-regulating equipment is required. Thus the flux-cored open arc hardfacing process is a simple semiautomatic welding process requiring only a torch and a device to feed the continuous electrode wire. Conventional power, supplied with drooping voltage-amperage characteristics, is well adapted to flux-cored open arc hardfacing equipment. Motor-generator units are frequently used for field applications because of their portability. In-plant operations normally make use of direct current rectifiers. Although flux-cored open arc hardfacing is a semiautomatic process, it is not necessary to have a constant-potential power supply for its operation. Flux-cored open arc equipment generally is designed for small-diameter welding wires, usually 1.2 mm ( (

3 in.), 1.6 mm 64

1 7 in.), or 2.8 mm ( in.). Most ferrous metals can be hardfaced with ease. Cast iron and chilled iron should be 16 64

handled cautiously. Open arc is a high-energy process capable of much higher deposition rates than SMAW. The 2.8 mm (

7 in.) diam rod can be deposited with welding currents as high as 500 A. The flux-cored open arc hardfacing process is 64

closely related to the submerged arc process, except that it does not require a flux burden over the bead; consequently,

them is little or no slag to be removed from the overlay. Most hardfacing alloys can be deposited by the flux-cored open arc process without preheat, and there is good recovery of elements across the arc. Hardfacing by fluxcored open arc welding is an easy process to use, and only a short time is required to train an operator for the average application. Arc stability and melting rates are excellent. The best hardfacing characteristics are obtained from DCEP. However, good results can be obtained with an alternating current supply at 26 to 30 V. Hardfacing deposits are comparable in soundness to deposits made with covered electrodes, at deposition rates three to five times faster than with covered electrodes. Hardfacing by flux-cored open arc welding is not without disadvantages. Because the arc is not shielded with inert gas or a blanket of granular flux, considerable spatter and some porosity can be expected with this process. In addition, because of high welding currents, it is not particularly well adapted to the hardfacing of small parts. The recovery of weld metal from the tube wire to the deposit normally averages near 80 to 85% or higher. The amount of cross cracking and distortion is comparable to that of other arc welding processes. Gas-metal arc welding is a consumable electrode hardfacing process in which the filler material and workpiece surfaces are protected by a flowing shielding gas--carbon dioxide, argon, or helium, either singly or in combination with a small amount of oxygen. The GMAW process is suitable for semiautomatic hardfacing, and the versatility of the process lends itself to the hardfacing of complex shapes. Hardfacing using GMAW can be fully mechanized. The hardfacing deposit is visible at all times in GMAW, thereby enabling high-quality deposits. Hardfacing by GMAW utilizes small-

diameter (generally 1.6 mm, or

1 in., or less) hardfacing filler wires or electrodes. Gas-metal arc hardfacing wires can be 16

deposited in either spray arc or short arc modes. The spray arc mode produces a continuous stream of droplets about the same diameter as the hardfacing wire, which are projected by electromagnetic force from the tip of the wire. Deposition rates are high, as is dilution. In the short arc mode, lower voltages are used, and transfer of the molten hardfacing wire is more globular, usually resulting in more spatter. Hardfacing by GMAW, unlike SMAW, requires a constant-potential power source with voltage, slope, and wire-feed rate controls. Best results for 1.6 mm (

1 in.) diam cobalt-base hardfacing wires in the spray arc mode are generally obtained 16

at 18 to 26 V, with wire-feed rates ranging from 45 to 65 mm/s (110 to 150 in./min.). Best results in the short arc mode are generally obtained at 16 to 18 V, with wire-feed rates ranging between 35 to 40 mm/s (80 to 100 in./min.). The use of auxiliary shielding gases adds to the cost of hardfacing by GMAW. However, the higher cost is generally offset by the higher deposit quality associated with GMAW. Gas tungsten arc welding is a nonconsumable electrode arc process in which the heated area of the workpiece, molten

hardfacing alloy, and the nonconsumable electrode are shielded from the atmosphere by a flow of shielding gas fed through the torch. Thoriated tungsten is the preferred nonconsumable electrode material. Conventional power supplies with drooping voltage-amperage characteristics are normally used for hardfacing. Motor-driven generators can be used, as can alternating current transformers with continuous high-frequency current. Direct current electrode negative is always used to minimize tungsten contamination in the hardfacing deposit. Argon generally is used as the shielding gas, but helium is also suitable. Hardfacing by manual GTAW is a useful alternative to hardfacing by OAW techniques, particularly when large components or reactive base metals are involved. Gas-tungsten arc welding is preferred for hardfacing on reactive base metals such as titanium-stabilized stainless steels or aluminum-bearing nickel-base alloys. Gas-tungsten arc welding also is preferred over OAW for hardfacing applications in which carbon pickup is unacceptable or in applications involving filler materials that tend to boil under the influence of an oxyacetylene flame. The localized, intense heat of GTAW generally results in more base metal dilution than OAW. Base metal dilution can be minimized by torch oscillation, by using no more amperage than is necessary, and/or by concentrating the arc on the hardfacing deposit rather than on the workpiece. Hardfacing deposition rates generally are comparable in manual GTAW and OAW, despite the difference in base metal dilution. Hardfacing by GTAW can be accomplished automatically by simply attaching the torch to an oscillating mechanism and using a mechanized device to feed the hardfacing filler material into the arc region. This, of course, requires long-length filler rod or continuous wire. High-quality, reproducible hardfacing deposits can be produced by automatic GTAW by controlling filler-metal feed rates, torch oscillation, and travel speed. Current-delay controls can be used to control the final stages of solidification in the hardfacing deposit to minimize shrinkage and crater cracking.

The arc action in GTAW is generally smoother, quieter, and freer of weld spatter than in other consumable electrode arc welding processes. The use of a separate filler rod or wire enables greater operator control of deposit shape. Gas-tungsten arc welding is more adaptable to hardfacing small intricate parts and generally produces higher-quality deposits than other hardfacing processes. Plasma Arc Welding. Hardfacing by PAW is similar to hardfacing by GTAW in that both processes employ a gas-

shielded arc between a nonconsumable tungsten electrode and the workpiece as the primary heat source for hardfacing. The processes differ in that the PAW process maximizes the use of plasma as a secondary heat source. The plasma is formed by ionizing the gas flowing in a nozzle surrounding the electrode. The electrode generally is recessed into the nozzle. The plasma gas generally emerges from a constricting orifice arrangement. The PAW process can use bare rod or wire as a hardfacing consumable, but more often powder is used as the consumable. When powder is used, the process is often referred to as the plasma transferred arc (PTA) process. Plasma transferred arc surfacing, as shown schematically in Fig. 10, is a welding process in which the powder is introduced into a combined arc/plasma stream to form a molten pool on the workpiece. The resulting deposit is homogeneous and dense, with excellent metallurgical bonding with the base metal (workpiece). This process uses a constricted arc. The alloy is carried from a powder feeder to the plasma torch in a stream of argon gas, which forms the plasma and is directed away from the torch into the arc effluent, where it is melted and fusion-bonded to the base metal. A direct current power source connecting the tungsten electrode and the workpiece provides energy for the transferred arc.

FIG. 10 SCHEMATIC OF THE PLASMA TRANSFERRED ARC HARDFACING PROCESS. SOURCE: REF 1

A second direct current power source connecting the tungsten electrode and the arc-constricting orifice supports a nontransferred arc. This nontransferred arc supplements the heat of the transferred arc and serves as a pilot arc to initiate the transferred arc. Argon is passed through a gas diffuser to provide a blanket of shielding gas in and around the arc zone. Hardfacing with the PTA process has many advantages. Deposits ranging from 0.64 to 3.2 mm (0.025 to 0.125 in.) thick can be produced rapidly in a single pass. A dilution of 5 to 20% in the deposit is typical. Powder recoveries as high as 95%, with deposition rates up to 4.5 kg/h (10 lb/h), are possible, depending on the size and shape of the part being hardfaced. Still higher deposition rates ranging up to 18 kg/h (40 lb/h) are possible using the plasma arc hot wire process, which, as its name implies, uses wire hardfacing consumables and two independently controlled systems operating in one weld pool to melt and fuse a hardfacing alloy to a substrate. This process, however, is considered a weld cladding operation and is described later in this article. Hardfacing by the PTA process has some disadvantages. The equipment is relatively expensive and is limited to straightline or cylindrical parts, except when special tooling is provided. Usually, several parts are needed for setup, which makes the process somewhat unattractive to the low-volume job shop. The torch is mounted on a carriage and track that require permanent installation. There are two circuits employed in this process; therefore, either a special power supply or two separate conventional power units are required. Argon consumption is somewhat higher than with GTAW. Argon has to

be supplied to the center of the torch to protect the electrode. Shielding gas is necessary to protect the weld metal, and argon is also used as a carrier for the powder. Large parts being faced with hard alloys normally require preheat. This is a serious potential problem with the PAW process, if excessive preheats and prolonged hardfacing times cause overheating of the torch. Laser Hardfacing (Ref 10) Laser hardfacing, also commonly referred to as laser cladding, differs little in principle from traditional forms of hardfacing; the primary difference is the use of a high-energy laser beam heat source rather than an arc or gas flame. Laser beams offer potential in applying thin overlays or when access to the surface to be hardfaced can be achieved more readily by a laser beam than with an electrode or torch. Cobalt-, nickel-, and tungsten carbide-base hardfacing alloys are the usual cladding materials used for laser hardfacing. As with conventional hardfacing methods, the materials are used in applications involving metal-to-metal contact, impact, erosion, and abrasion wear resistance. Other laser hardfacing materials include titanium carbide, Fe-Cr-Ni-B alloys, aluminum bronzes, and ceramics. Substrates have included carbon and low-alloy steels, stainless steels, nickel-base alloys (including superalloys), aluminum, cast irons, and tool steels (Ref 11). Hardfacing material selection depends on its metallurgical compatibility with the substrate material. Alloys that form brittle intermetallic phases with the substrate are undesirable. For example, cobalt forms harmful intermetallics with titanium, and cobalt-base hardfacing alloys would be unsuitable on titanium alloy substrates. Processing. The hardfacing alloy is melted by a laser beam and allowed to spread freely and freeze over the substrate.

The beam also melts a very thin layer of the substrate, which combines with the liquid weld metal to the least extent necessary and solidifies to form a strong metallurgical bond. A good fusion bond can be achieved with a dilution zone that is only 10 to 20 μm thick. The hardfacing alloy can be in several forms, examples of which are a prealloyed powder that is applied to the sample surface with or without a binder, a self-fluxing powder that is flame-sprayed, a hardfacing alloy that is plasma-sprayed, or a chip that is preplaced. Laser consolidation of these coatings results in densification and smoothening, eliminates channels to the substrate, improves the bonding between coating and substrate, and reduces porosity, all of which contribute to the strength and integrity of the hardfacing layer. Processing parameters for laser cladding are a power density that ranges from 10 to 1000 MW/m2 and an interaction time from 0.1 to 1 s. The shielding gas could be any of the inert gases or a combination of gases, such as He/Ar and H2/Ar. To date the most successful application of laser hardfacing involves a proprietary process that utilizes a specially designed powder-feed apparatus. Using such equipment, the powder and an assist gas are fed to the weld area through a ceramic nozzle; a shielding gas of helium and argon surrounds the powder-gas mixture as it leaves the nozzle (Fig. 11). The powder delivery nozzle is positioned in such a manner as to flood the entire melt pool with powder. The feeding angle is generally from 35 to 45° to the horizontal, and the feed tube, which typically has a 3 mm (0.12 in.) diam, is positioned 10 to 12 mm (0.4 to 0.5 in.) from the substrate. Typically, the powder flow rate is from 0.005 to 0.1 cm3/s (0.0003 to 0.006 in.3/s), particle velocity is from 1 to 2 m/s (3.3 to 6.6 ft/s), and carrier gas velocity is from 3 to 7 m/s (10 to 23 ft/s). Desirable dilutions are from 3 to 8%. Overlay thickness can be varied from 0.15 to 4 mm (0.005 to 0.15 in.). Uniform feeding of the powder ensures uniform surfacing layers. Unlike plasma-sprayed coatings, porosity and unmelted powder particles are almost never observed within a laser hardfacing layer.

FIG. 11 SCHEMATIC OF THE LASER HARDFACING PROCESS USING DYNAMIC POWDER FEED. COURTESY OF QUANTUM LASER CORPORATION

Laser hardfacing deposits are commonly achieved by relatively large doughnut-shape beams, or by somewhat focused beams that are oscillated. These ensure a desirable deposit profile. Hardfacing material can also be added in rod, wire, or sheet form, but special procedures are needed because of reflectivity problems. Thermal stresses in the weld metal can cause harmful cracking, but this can be eliminated by an appropriate preheating practice. Low power densities, large beam diameters, and slow sample translation rates tend to produce crack-free deposits. Laser Hardfacing Applications. Because one of the primary goals of laser hardfacing is to improve the tribological properties of components, several applications have been found for the technique. Some of them are in the exploratory stage, others are in the pilot-plant stage, and a few have reached the production stage.

Ayers and Gnanamuthu (Ref 12) give several examples of laser hardfacing for improved wear resistance, such as Tribaloy T-800 on powder-coated ASTM A 387 steel, Haynes Stellite alloy No. 1 cast rods placed on AISI 4815 steel, and a mixture of WC and iron powder applied without binder on AISI 1018 steel. In addition to stainless steel on mild steel, Steen (Ref 13) has reported successful laser hardfacing of nickel, bronze, Fe-B, and Stellite SF6 on mild steel; chromium on titanium; stainless on aluminum; mild steel on stainless; and Stellite on brass. Other examples of laser hardfacing are AISI 1016 steel clad with a mixture of chromium, manganese, and carbon powders (Ref 14) and gray cast iron hardfaced with a powder-fed mixture of molybdenum, chromium-carbide (Cr3C2), chromium, nickel, and silicon (Ref 15). Bruck (Ref 16) demonstrated laser hardfacing within a confined space by coating the inside surface of small-bore pipes (inside diameter of 50 to 100 mm, or 2 to 4 in., and length of 0.3 to 1.2 m, or 1 to 4 ft). An oscillating mirror placed in the pipe formed a 12.5 mm (0.5 in.) wide melt pass, into which the hardfacing alloy powder was fed. The pipe was rotated and translated to cover the entire inner surface. Previously, chromium plating was used, but the laser technique appreciably improved galling resistance. Using suitable cooling techniques, the pipe temperature was maintained below 500 °C (930 °F) to avoid deterioration of the core properties, as well as distortion. La Rocca (Ref 17) describes laser hardfacing of exhaust valves with a cobalt-base alloy at an Italian automotive manufacturer. The cobalt-base powder was fed into the laser melt pool. The laser hardfacing material was superior to GTAW material in terms of thickness uniformity, uniform microstructure and elemental distributions, and adhesion. When compared to GTAW, excess-metal removal was reduced by 10 to 15%. One British automotive manufacturer used production-stage laser hardfacing of a nickel-alloy turbine blade shroud interlock by powder feeding either a Tribaloy

alloy or a Nimonic alloy (Ref 18). Previously, a manual arc melting technique was used. Laser hardfacing reduced hardfacing time from 14 min to 75 s, improved productivity and quality, reduced cost by 85%, and reduced powder consumption by 50%. Table 11 lists components that are laser hardfaced for commercial applications.

TABLE 11 TYPICAL APPLICATIONS FOR LASER HARDFACING

COMPONENT GATE AND SEAT OF STEEL VALVES FOR OIL-FIELD, GEOTHERMAL, AND NUCLEAR ENERGY PRODUCTION, AND CHEMICAL PROCESSING DIESEL ENGINE VALVE

HARDFACING ALLOY STELLITE 6

HARDFACING TECHNIQUE POWDER FEED

STELLITE SF6

GATE VALVE SEAT

DELORO 60

GAS TURBINE BLADE "Z" NOTCH LEADING EDGE OF TURBINE BLADES PUMP BUSHINGS AND IMPELLERS ENGINE VALVE BASES

CO-BASE ALLOY

PREPLACED POWDER PASTE PREPLACED POWDER PASTE POWDER FEED

CO-BASE ALLOY

POWDER FEED

CO-BASE ALLOY

POWDER FEED

NI-CR ALLOY

POWDER FEED

OIL-FIELD VALVE GATES

WC AND CO ALLOY

POWDER FEED

TRACTOR BUSHINGS

STELLITE 6

POWDER FEED

TURBINE BLADES, PUMP VALVE SEATS

STELLITE, COLMOLOY, BLENDED POWDER CRC2, CR, NI, MO/CAST FE STELLITE, COLMONOY, ALLOYS/CARBIDES

PREPLACED BEDS, GRAVITY FEED PREPLACED POWDER POWDER FEED

STELLITE, TRIBALOY

POWDER FEED

PWA 694, NIMONICS

PREPLACED CHIP

VALVE STEM, VALVE SEAT, ALUMINUM BLOCK OFF-SHORE DRILLING AND PRODUCTION PARTS, VALVE COMPONENTS, BOILER FIREWALL AEROSPACE COMPONENTS TURBINE BLADE SHROUD INTERLOCKS

COMPANY W-K-M DIV. OF JOY MFG., HOUSTON, TX

CABOT CORP., KOKOMO, IN CABOT CORP., KOKOMO, IN QUANTUM LASER, EDISON, NJ QUANTUM LASER, EDISON, NJ QUANTUM LASER, EDISON, NJ QUANTUM LASER, EDISON, NJ QUANTUM LASER, EDISON, NJ QUANTUM LASER, EDISON, NJ WESTINGHOUSE, PITTSBURGH, PA FIAT, TURIN, ITALY COMBUSTION ENG., OH

ROCKWELL INT., CA PRATT & WHITNEY, WEST PALM BEACH, FL

Source: Ref 10


1. K.C. ANTONY ET AL., HARDFACING, WELDING, BRAZING, AND SOLDERING, 9TH ED., VOL 6, METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 771-793 2. P. CROOK AND H.N. FARMER, FRICTION AND WEAR OF HARDFACING ALLOYS, FRICTION,

LUBRICATION, AND WEAR TECHNOLOGY, VOL 18, ASM HANDBOOK, ASM INTERNATIONAL, 1992, P 758-765 3. D.K. SUBRAMANYAM, A.E. SWANSIGER, AND H.S. AVERY, AUSTENITIC MANGANESE STEELS, PROPERTIES AND SELECTION: IRONS, STEELS, AND HIGH-PERFORMANCE ALLOYS, VOL 1, ASM HANDBOOK, ASM INTERNATIONAL, 1990, P 822-840 4. H.S. AVERY AND H.J. CHAPIN, WELD J., VOL 31, 1952, P 917 5. R.D. ZORDAN, REPORT 10890, CABOT CORPORATION, KOKOMO, IN, 1982 6. H. OCKEN AND A.K. VELAN, ED., "LABORATORY EVALUATIONS OF COBALT-FREE, NICKELBASED HARDFACING ALLOYS FOR NUCLEAR APPLICATIONS," REPORT NP-4993, PROJECT 1935, ELECTRIC POWER RESEARCH INSTITUTE, PALO ALTO, CA, MARCH 1987 7. P.J. HOFMANN, B.C. FRIEDRICH, AND H. OCKEN, ED., "LABORATORY EVALUATIONS OF IRON-BASED HARDFACING ALLOYS," REPORT NP-5874, PROJECT 1935-11, ELECTRIC POWER RESEARCH INSTITUTE, PALO ALTO, CA, JUNE 1988 8. WELD DESIGN AND FABRICATION, DATA SHEET 420, WELD. DES. FABR., VOL 48 (NO. 8), 1975, P 50 9. K.M. KULKARNI AND V. ANAND, METAL POWDERS USED FOR HARDFACING, POWDER METALLURGY, 9TH ED., VOL 7, ASM HANDBOOK, AMERICAN SOCIETY FOR METALS, 1984, P 823-836 10. K.P. COOPER, LASER SURFACE PROCESSING, FRICTION, LUBRICATION, AND WEAR TECHNOLOGY, VOL 18, ASM HANDBOOK, ASM INTERNATIONAL, 1992, P 861-872 11. A.G. BLAKE, ET AL., "LASER COATING TECHNOLOGY: A COMMERCIAL REALITY," BULLETIN OF QUANTUM LASER CORPORATION, EDISON, NJ 12. J.D. AYERS AND D.S. GNANAMUTHU, HARDENING OF METAL SURFACES BY LASER PROCESSING, WELDING, BRAZING, AND SOLDERING, 9TH ED., VOL 6, ASM HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 793-803 13. W.M. STEEN, PROC. NATO ADVANCED STUDY INSTITUTE ON LASER SURFACE TREATMENT OF METALS, C.W. DRAPER AND P. MAZZOLDI, ED., MARTINUS NIJHOFF, 1986, P 369 14. J. MAZUMDER AND J. SINGH, PROC. NATO ADVANCED STUDY INSTITUTE ON LASER SURFACE TREATMENT OF METALS, C.W. DRAPER AND P. MAZZOLDI, ED., MARTINUS NIJHOFF, 1986, P 297 15. A. BELMONDO AND M. CASTAGNA, SOURCEBOOK ON APPLICATIONS OF THE LASER IN METALWORKING, E. METZBOWER, ED., AMERICAN SOCIETY FOR METALS, 1981, P 310 16. G.J. BRUCK, J. MET., VOL 39 (NO. 2), 1987, P 10 17. A.V. LA ROCCA, PROC. NATO ADVANCED STUDY INSTITUTE ON LASER SURFACE TREATMENT OF METALS, C.W. DRAPER AND P. MAZZOLDI, ED., MARTINUS NIJHOFF, 1986, P 521 18. R.M. MACINTYRE, PROC. NATO ADVANCED STUDY INSTITUTE ON LASER SURFACE TREATMENT OF METALS, C.W. DRAPER AND P. MAZZOLDI, ED., MARTINUS NIJHOFF, 1986, P 545 Hardfacing, Weld Cladding, and Dissimilar Metal Joining J.R. Davis, Davis & Associates

Thermal Spraying Thermal spraying comprises a group of processes in which divided molten metallic or nonmetallic material is sprayed onto a prepared substrate to form a coating. The sprayed material is originally in the form of wire, rod, or powder. As the coating materials are fed through the spray unit, they are heated to a molten or plastic state and propelled by a stream of compressed gas onto the substrate. As the particles strike the surface, they flatten and form thin platelets that conform and

adhere to the irregularities of the prepared surface and to each other. They cool and accumulate, particle by particle, into a lamellar, castlike structure. In general, the substrate temperature can be kept below approximately 200 °C (400 °F), eliminating metallurgical change of the substrate material. The spray gun generates the necessary heat for melting through combustion of gases, an electric arc, or a plasma. Figure 12 illustrates a general thermal spraying process.

FIG. 12 SCHEMATIC OF THE GENERAL THERMAL SPRAY PROCESS. SOURCE: REF 20

The deposited structures of thermal spray coatings differ from those of the same material in the wrought form because of the incremental nature of the coating buildup and because the coating composition is often affected by reaction with the process gases and the surrounding atmosphere while the materials are in the molten state. For example, where air or oxygen is used as the process gas, oxides of the material applied may be formed and become part of the coating. The asapplied structures of all thermal spray coatings are similar in their lamellar nature; the variations in structure depend on the particular thermal spray process used, the processing parameters and techniques employed, and the material applied. Figure 13 illustrates the microstructure that results from the thermal spray process. As shown in this figure, the molten particles spread out and deform (splatter) as they impact the substrate, at first locking onto irregularities on the roughened surface, then interlocking with each other.

FIG. 13 SCHEMATIC SHOWING THE BUILDUP OF A THERMAL SPRAY COATING. MOLTEN PARTICLES SPREAD OUT AND DEFORM (SPLATTER) AS THEY STRIKE THE TARGET, AT FIRST LOCKING ONTO IRREGULARITIES ON THE SUBSTRATE, THEN INTERLOCKING WITH EACH OTHER. VOIDS CAN OCCUR IF THE GROWING DEPOSIT TRAPS AIR. PARTICLES OVERHEATED IN THE FLAME BECOME OXIDIZED. UNMELTED PARTICLES MAY SIMPLY BE EMBEDDED IN THE ACCUMULATING DEPOSIT. COURTESY OF THE MATERIALS ENGINEERING INSTITUTE, ASM INTERNATIONAL

The bond between the sprayed coating and the substrate is generally mechanical. Proper surface preparation of the substrate before spraying is often the most critical influence on the bond strength of the coating. Surface preparation techniques will be discussed later in this section. Thermal spraying is a very versatile technique in that it can be used both in-shop or on-site. The thermal spray process permits rapid application of high-performance material in thicknesses from a few mils to more than 25 mm (1 in.) thick on parts of a variety of sizes and geometries. A list of applications would include coatings for: • • • • • • •

WEAR RESISTANCE OXIDATION RESISTANCE CORROSION RESISTANCE DIMENSIONAL RESTORATION THERMAL BARRIERS ELECTRICAL CONDUCTIVITY OR RESISTIVITY BIOMEDICAL DEVICES

Thermal Spray Processes (Ref 19, 20) Currently, five different commercially available thermal spray methods are in use: • • • • •

OXYFUEL WIRE (OFW) SPRAY ELECTRIC ARC WIRE (EAW) SPRAY OXYFUEL POWDER (OFP) SPRAY PLASMA ARC (PA) POWDER SPRAY HIGH-VELOCITY OXYFUEL (HVOF) POWDER SPRAY

Of these, the plasma arc process is most commonly used. Market studies indicate that by the year 2000, plasma spraying will capture about 48% of the market, the HVOF process 25%, the electric-arc method 15%, and the combustion methods (OFW and OFP) a combined 12% (Ref 21). Table 12 compares important process characteristics associated with these techniques. Selection of the appropriate thermal spray method is typically determined by: • • • •

DESIRED COATING MATERIAL COATING PERFORMANCE REQUIREMENTS ECONOMICS PART SIZE AND PORTABILITY

TABLE 12 COMPARISON OF THERMAL SPRAY PROCESSES PROCESS

FLAME POWDER: POWDER FEEDSTOCK, ASPIRATED INTO THE OXYGEN/FUEL-GAS FLAME, IS MELTED AND CARRIED BY THE FLAME ONTO THE WORKPIECE. PARTICLE VELOCITY IS RELATIVELY LOW, AND BOND STRENGTH OF DEPOSITS IS LOW. POROSITY IS HIGH AND COHESIVE STRENGTH IS LOW. SPRAY RATES ARE USUALLY IN THE 0.5 TO 9 KG/H (1 TO 20 LB/H) RANGE. SURFACE TEMPERATURES CAN RUN QUITE HIGH. FLAME WIRE: IN FLAME WIRE SPRAYING, THE ONLY FUNCTION OF THE FLAME IS TO MELT THE MATERIAL. A STREAM OF AIR THEN DISINTEGRATES THE MOLTEN MATERIAL AND PROPELS IT ONTO THE WORKPIECE. SPRAY RATES FOR MATERIALS SUCH

GAS FLOW M3/H 11

FT3/H 400

FLAME OR EXIT PLASMA TEMPERATURE °C °F 2200 4000

71

2500

2800

5000

PARTICLE IMPACT VELOCITY M/S FT/S 30 100

RELATIVE ADHESIVE STRENGTH(A)

COHESIVE STRENGTH

OXIDE CONTENT, %

RELATIVE PROCESS COST(A)

3

LOW

6

3

MAXIMUM SPRAY RATE KG/H LB/H 7 15

180

4

MEDIUM

4

3

9

600

20

POWER

KW 2575

HP 34100

50100

70135

ENERGY REQUIRED TO MELT KW/KG KW/LB 11-22 5-10

11-22

5-10

AS STAINLESS STEEL ARE IN THE RANGE OF 0.5 TO 9 KG/H (1 TO 20 LB/H). SUBSTRATE TEMPERATURES ARE FROM 95 TO 205 °C (200 TO 400 °F) BECAUSE OF THE EXCESS ENERGY INPUT REQUIRED FOR FLAME MELTING. WIRE ARC: TWO CONSUMABLE WIRE ELECTRODES ARE FED INTO THE GUN, WHERE THEY MEET AND FORM AN ARC IN AN ATOMIZING AIR STREAM. THE AIR FLOWING ACROSS THE ARC/WIRE ZONE STRIPS OFF THE MOLTEN METAL, FORMING A HIGH-VELOCITY SPRAY STREAM. THE PROCESS IS ENERGY EFFICIENT; ALL INPUT ENERGY IS USED TO MELT THE METAL. SPRAY RATE IS ABOUT 2.3 KG/H/KW (5 LB/H/KW). SUBSTRATE TEMPERATURE CAN BE LOW BECAUSE ENERGY INPUT PER POUND OF METAL IS ONLY ABOUT ONEEIGHTH THAT OF OTHER SPRAY METHODS. CONVENTIONAL PLASMA: CONVENTIONAL

71

2500

5500

10,000

240

800

6

HIGH

0.5-3

1

16

35

4-6

5-8

0.2-0.4

0.1-0.2

4.2

150

5500

10,000

240

800

6

HIGH

0.5-1

5

5

10

3080

40110

13-22

6-10

PLASMA SPRAYING PROVIDES FREEPLASMA TEMPERATURES IN THE POWDER HEATING REGION OF 5500 °C (10,000 °F) WITH ARGON PLASMA, AND 4400 °C (8000 °F) WITH NITROGEN PLASMA--ABOVE THE MELTING POINT OF ANY KNOWN MATERIAL. TO GENERATE THE PLASMA, AN INERT GAS IS SUPERHEATED BY PASSING IT THROUGH A DC ARC. POWDER FEEDSTOCK IS INTRODUCED AND IS CARRIED TO THE WORKPIECE BY THE PLASMA STREAM. PROVISIONS FOR COOLING OR REGULATION OF THE SPRAY RATE MAY BE REQUIRED TO MAINTAIN SUBSTRATE TEMPERATURES IN THE 95 TO 205 °C (200 TO 400 °F) RANGE. TYPICAL SPRAY RATE IS 0.1 KG/H/KW (0.2 LB/H/KW). DETONATION GUN: SUSPENDED POWDER IS FED INTO A 1 M (3 FT) LONG TUBE ALONG WITH OXYGEN AND FUEL GAS. A

11

400

3900

7000

910

3000

8

VERY HIGH

0.1

10

1

2

100270

135360

220

100

SPARK IGNITES THE MIXTURE AND PRODUCES A CONTROLLED EXPLOSION. THE HIGH TEMPERATURES AND PRESSURES (1 MPA, OR 150 PSI) THAT ARE GENERATED BLAST THE PARTICLES OUT OF THE END OF THE TUBE TOWARD THE SUBSTRATE. HIGH-VELOCITY OXYFUEL: IN HVOF SPRAYING, A FUEL GAS AND OXYGEN ARE USED TO CREATE A COMBUSTION FLAME AT 2500 TO 3100 °C (4500 TO 5600 °F). THE COMBUSTION TAKES PLACE AT VERY HIGH CHAMBER PRESSURE, EXITING THROUGH A SMALLDIAMETER BARREL TO PRODUCE A SUPERSONIC GAS STREAM AND VERY HIGH PARTICLE VELOCITIES. THE PROCESS RESULTS IN EXTREMELY DENSE, WELLBONDED COATINGS, MAKING IT ATTRACTIVE FOR MANY CORROSIONRESISTANT APPLICATIONS.

28-57

10002000

3100

5600

6101060

20003500

8

VERY HIGH

0.2

5

14

30

100270

135360

22-200

10-90

EITHER POWDER OR WIRE FEEDSTOCK CAN BE SPRAYED, AT TYPICAL RATES OF 2.3 TO 14 KG/H (5 TO 30 LB/H). HIGH-ENERGY PLASMA: THE HIGHENERGY PLASMA PROCESS PROVIDES SIGNIFICANTLY HIGHER GAS ENTHALPIES AND TEMPERATURES, ESPECIALLY IN THE POWDER HEATING REGION, DUE TO A MORE STABLE, LONGER ARC AND HIGHER POWER DENSITY IN THE ANODE NOZZLE. THE ADDED POWER (TWO TO THREE TIMES THAT OF CONVENTIONAL PLASMA) AND GAS FLOW (TWICE AS HIGH) PROVIDE A LARGER, HIGHER TEMPERATURE POWDER INJECTION REGION AND REDUCED AIR ENTRAINMENT. ALL THIS LEADS TO IMPROVED POWDER MELTING, FEWER UNMELTS, AND HIGH PARTICLE IMPACT VELOCITY. VACUUM PLASMA: VACUUM PLASMA USES A CONVENTIONAL PLASMA TORCH IN A CHAMBER AT

17-28

6001000

8300

15,000

2401220

8004000

8

VERY HIGH

0.1

4

23

50

100250

135335

9-13

4-6

8.4

300

8300

15,000

240610

8002000

9

VERY HIGH

(B)

10

10

24

50100

70135

11-22

5-10

PRESSURES IN THE RANGE OF 10 TO 50 KPA (0.1 TO 0.5 ATM). AT LOW PRESSURES THE PLASMA IS LARGER IN DIAMETER, IS LONGER, AND HAS A HIGHER VELOCITY. THE ABSENCE OF OXYGEN AND THE ABILITY TO OPERATE WITH HIGHER SUBSTRATE TEMPERATURES PRODUCE DENSER, MORE ADHERENT COATINGS HAVING MUCH LOWER OXIDE CONTENTS. Source: Ref 21

(A) (B)

1 (LOW) TO 10 (HIGH). PPM LEVELS.

The oxyfuel wire spray process (also called wire flame spraying or the combustion wire process) is the oldest of the

thermal spray coating methods and among the lowest in capital investment. The process utilizes an oxygen-fuel gas flame as a heating source and coating material in wire form. Any material in the form of wire and capable of being melted below 2480 °C (4500 °F) can be flame-sprayed. Solid-rod feed stock has also been used. During operation, the wire is drawn into the flame by drive rolls that are powered by an adjustable air turbine or electric motor. The tip of the wire is melted as it enters the flame, atomized into particles by a surrounding jet of compressed air, and propelled to the workpiece. Figure 14 shows a schematic of the OFW spray process.

FIG. 14 SCHEMATIC OF THE OXYFUEL WIRE SPRAY PROCESS. SOURCE: REF 20

Spray rates for this process range from 0.5 to 10 kg/h (1 to 20 lb/h) and are dictated by the melting point of the material and the choice of fuel gas. Common fuel gases are acetylene, MAPP gas, propane, propylene, and natural gas, each combined with oxygen. The wire spray gun is most commonly used as a hand-held device for on-site application, although an electric-motor-driven gun is recommended for fixed-mounted use in high-volume, repetitive production work. The OFW process is widely used for corrosion protection of large outdoor structures, such as bridges and storage tanks, and for restoration of dimension to worn machinery components. It is a good choice for all-purpose spraying. Coatings can be applied rapidly and at low cost, and a wide variety of metal coating materials are available. Typical spray materials include austenitic and martensitic stainless steels, nickel aluminide, nickel-chromium alloys, bronze, Monel, babbitt, aluminum, zinc, and molybdenum. In general, oxyfuel wire flame-sprayed coatings exhibit lower bond strengths, higher porosity, a narrower working temperature range, and higher heat transmittal to the substrate than plasma or electric arcsprayed coatings. Electric arc wire spraying also applies coatings of selected metals in wire form. Push-pull motors feed two electrically charged wires through the arc gun to contact tips at the gun head. An arc is created that melts the wires at temperatures above 5500 °C (10,000 °F). Compressed air atomizes the molten metal and projects it onto a prepared surface. Figure 15 shows a schematic of the EAW spray process.

FIG. 15 SCHEMATIC OF THE ELECTRIC ARC WIRE SPRAY PROCESS. SOURCE: REF 20

The EAW process is excellent for applications that require heavy coating buildup or have large surfaces to be sprayed. The arc system can produce a spray pattern ranging from 50 to 300 mm (2 to 12 in.) and can spray at high speeds. It has built-in flexibility, allowing coating characteristics, such as hardness or surface texture, to be tailored to specific applications. The EAW method is characterized by strong coating adhesion because of the high particle temperatures produced. Because the process uses only electricity and compressed air, it allows equipment to be moved relatively easily from one installation to another, and it eliminates the need to stock oxygen and fuel gas supplies. Materials applied by the EAW method are similar to those used in the OFW process. The oxyfuel powder spray method, which also uses combustion gases as the heat source, extends the range of available coatings and subsequent applications to include ceramics, cermets, carbides, and fusible hardfacing coatings. Using either gravity flow or pressurized feed, powder is fed into the gun and carried to the gun nozzle, where it is melted and projected by the gas stream onto a prepared surface (Fig. 16). For general-purpose spraying, the gravity-flow system is used. When exacting coating consistency and/or high spray rates are desired, the pressurized feed system is used. Oxyfuel powder guns are the lowest-cost thermal spray equipment, are easiest to set up, and make it easy to change coating materials. The OFP method finds widest use in short-run machinery maintenance work and in the production spraying of abradable clearance-control seals for gas turbine engines. The OFP spray equipment is lighter and more compact than other forms of thermal spray equipment, but because of the lower particle velocities obtained, the coatings generally have lower bond strength, higher porosity, and a lower overall cohesive (interparticle) strength than the coatings produced by other thermal spray processes.

FIG. 16 SCHEMATIC OF THE OXYFUEL POWDER SPRAY PROCESS. SOURCE: REF 20

Flame spray and fuse is a modification of the OFP spray method. The materials used for coating are self-fluxing, fusible materials which require postspray heat treatment. In general, these are the boride-containing nickel-base alloys described earlier in this article. In practice, parts are prepared and coated as in other thermal spray processes and then fused. Fusing is accomplished using one of several techniques, such as flame or torch; induction; or in vacuum, inert, or hydrogen furnaces. These alloys generally fuse between 1010 to 1175 °C (1850 to 2150 °F), depending upon composition. Spray-and-fuse coatings are used in applications where excessive wear is a problem. Coating hardnesses can be as high as 65 HRC. The use of spray-and-fuse coatings is limited to substrate materials that can tolerate the 1010 to 1175 °C (1850 to 2150 °F) fusing temperatures. Fusing temperatures may alter the heat-treated properties of some alloys. However, the coating will usually withstand additional heat treatment of the substrate. Plasma arc spraying is one of the most sophisticated and versatile thermal spray methods. The PA spray process uses

coating materials in the form of powders, which are heated with a plasma heat source (Fig. 17). Temperatures that can be obtained with commercial plasma equipment have been calculated to be greater than 11,000 °C (20,000 °F) and are far above the melting or even the vaporization point of any known material. Decomposition of materials during spraying is minimized because of the high gas velocities produced by the plasma, resulting in extremely short residence time in the thermal environment. The plasma process also provides a controlled atmosphere for melting and transport of the coating material, thus minimizing oxidation, and the high gas velocities produce coatings of high density.

FIG. 17 SCHEMATIC OF THE PLASMA ARC SPRAY PROCESS. COURTESY OF D. BAXTER, ADVANCED MATERIALS & PROCESSES

The plasma gun operates on the principle of raising the energy state of a gas by passing it through an electric arc. The release of energy in returning the gas to its ground state results in exceedingly high temperatures. A gas such as nitrogen or argon enters a direct-current arc between a tungsten cathode and a copper anode that make up the nozzle (Fig. 17). Both components are cooled by a constant flow of water through internal passages. Here the plasma gas first dissociates (in the case of nitrogen, into two atoms), followed by ionization that releases free electrons. The electrons recombine outside the electric arc, and energy is released as heat and light. In addition, frequent collisions transfer energy from the electrons to the positive ions, accelerating them until the plasma reaches a state of equilibrium. The result is a thermal plasma in which the energy of the electrons has been turned into enthalpy, or heat content. At this point, powdered coating material suspended in a gas is injected into the plasma and is subsequently melted and propelled at high velocity to the workpiece. In practice, a small amount of a secondary gas, such as hydrogen or helium, is mixed with the primary plasma gas to increase operating voltage and thermal energy. The high temperatures and high gas velocities produced by the plasma process result in coatings that are superior in mechanical and metallurgical properties to low-velocity OFW or OFP coatings. The plasma process is particularly efficient for spraying high-quality coatings of ceramic materials, such as zirconium oxide for turbine engine combustors and chromium oxide for printing rolls. The plasma process is also readily field-portable and is thus used for large on-site applications such as power utility plant boiler tubes. Current plasma spray technology permits fully automatic start/stop operation and closed-loop computer control for power level, plasma gas flow, and powder feed rate. System problems can be diagnosed via computer modem to the equipment manufacturer, and system performance can be documented and stored on a digital recorder or data logger.

With any thermal spray powder method, the degree to which a given flame effectively melts and accelerates the powder depends on the type of coating material and the size and shape of the particles. Each particular coating material and gun combination has an optimum particle size. Particles much smaller than ideal will overheat and vaporize; much larger particles will not melt and may fall from the flame or rebound from the target. For the highest coating density, powders are sized between about 44 μm (325 mesh) and 10 μm (Ref 1, 9). The lower limit ensures free flow. Coarser powder may be sized up to 100 μm, with typical distribution falling between 88 μm (170 mesh) and 44 μm (325 mesh). Powder shape ranges from spherical to acicular; spherical is best because of reduced surface area and consistent flow. Any material that melts without subliming and is available as a properly sized powder has potential as a plasma spray process coating. Plasma spray coatings may be elemental (aluminum, molybdenum, nickel, or chromium), alloyed (nickel-, iron-, or cobalt-base), ceramics and cermets (Al2O3, Cr2O3, or Cr3C2), composites (nickel-aluminum, cobaltbonded tungsten carbide, or clad Al2O3), or mechanical blends. In one process variation, the plasma gun and workpiece are enclosed in a vacuum chamber, and the entire operation is carried out in an inert atmosphere at low pressure (7 kPa, or 50 torr). Advantages of this low-pressure plasma technique include improved bonding and density of the deposit, improved control over coating thickness (even with an irregular work surface), and higher deposit efficiency. However, facility cost is higher, and workpiece size is limited by the vacuum chamber size. This modification is suitable for coating turbine airfoils, blade tips, and shroud segments with complex coatings such as NiCrAlY and other MCrAlY-type alloys. Very fine powder (-325 mesh) generally is used for low-pressure plasma spraying to achieve maximum density. If high-purity coatings are also desired, then the powder must be manufactured under inert cover and atomized in an inert gas to preserve purity. The high-velocity oxyfuel powder spray process represents the state of the art for thermal spray coatings. The

HVOF process uses extremely high kinetic energy and controlled thermal energy output to produce very-low-porosity coatings that exhibit high bond strength, fine as-sprayed surface finish, and low residual stresses. Both pulsed-combustion and continuous-combustion HVOF methods are used. Pulsed-Combustion HVOF. The detonation gun (or D-gun) process was the first HVOF method introduced and was

initially developed for the deposition of hard, wear-resistant materials such as oxides and carbides (Ref 1, 9). The extremely high particle velocities achieved in the detonation gun result in coatings with higher density, greater internal strength, and superior bond strength than can be achieved with conventional plasma spraying or oxyfuel powder spraying. Detonation gun coatings have been successfully applied to critical areas of precision components made from virtually all commercial alloys. The typical coating thickness is less than 0.75 mm (0.030 in.), but many applications require only 0.075 mm (0.003 in.) or less of finished coating. This process, which is depicted schematically in Fig. 18, uses a controlled series of explosions of a mixture of oxygen and acetylene to blast the powder onto the surface of the workpiece. The detonation gun consists of a water-cooled barrel several feet long, with an inside diameter of about 25 mm (1 in.), and associated gas- and powder-metering equipment. In operation, a mixture of oxygen and acetylene is fed into the barrel along with a charge of powder. The gas is then ignited, and the detonation wave accelerates the powder to about 760 m/s (2500 ft/s), while heating it close to, or above, its melting point. Temperatures up to 3900 °C (7000 °F) are reached inside the gun. The distance that the powder is entrained in the high-velocity gas is much longer than in a plasma or flame spray device, which accounts for the high particle velocity. After the powder has exited the barrel, a pulse of nitrogen purges the barrel. The gun discharges four to eight times per second.

FIG. 18 SCHEMATIC OF THE DETONATION GUN (D-GUN) HVOF PROCESS. SOURCE: REF 1

Suitable materials for detonation gun coatings include pure metals, alloys, oxides, carbides, composites, and mechanical blends of two or more components. Historically, the detonation gun has proven particularly advantageous in applying hard wear-resistant carbide and oxide coatings. Detonation gun coatings have been successfully applied to critical areas of precision parts made from virtually all commercial alloys. The substrate surface is seldom heated above 150 °C (300 °F) during the coating process. As a result, a part can be fabricated and fully heat treated prior to coating. In addition, very high melting point materials (such as WC-Co) can be deposited on relatively low melting point substrates (such as aluminum). Usually, there is little or no distortion, and the substrate microstructure and strength are not changed by the coating process. Continuous-combustion HVOF processes, which also feature internal, confined combustion, operate with an oxyfuel mixture consisting of oxygen and either propylene, propane, hydrogen, and even liquid fuels such as kerosene. Highvelocity oxyfuel guns have combustor sections (air- or water-cooled), which combust the fuel/oxygen mixtures under pressure and accelerates the combusted gas streams down a confined cooled tube. Powders, with carrier gases, are fed into the nozzle where the particles are entrained with the confined, high-pressure combusting gases.

The high-velocity gases produce uniquely characteristic multiple shock diamond patterns, which are visible in the flame (Fig. 19). Combustion temperatures approach 2750 °C (5000 °F) and form a circular flame configuration. Powder is injected into the flame axially to provide uniform heating, and powder particles are accelerated by the high velocity gases, which typically approach speeds ranging from 1370 m/s (4500 ft/s) to as high as 2160 m/s (7100 ft/s). Particle velocities range from 760 m/s (2500 ft/s) to more than 900 m/s (3000 ft/s).

FIG. 19 SCHEMATICS OF (A) AIR-COOLED AND (B) WATER-COOLED CONTINUOUS COMBUSTION HVOF THERMAL SPRAY PROCESSES. (A) COURTESY OF METCO PERKIN ELMER. (B) SOURCE: REF 22

Continuous-combustion HVOF spray coating processes are separated into four major designs: (1) internal combustion, with the hot gases directed into a chamber where they mix with and heat a mixture of carrier gas and powder particles; (2) simultaneous injection of powder into a combustion chamber, where the mixture continuously bums and heats the particles; (3) confined combustion into which a stream of powder with carrier gas is injected; and (4) high-pressure liquid fuel systems with radial injection of powders downstream of the combustor exit. Figure 19 illustrates two of these variations. Common to all HVOF methods is that the gases are being burned under higher-than-atmospheric pressures, with the combustion gases, carrier gas, and an entrained distribution of powders being accelerated in the tube. The length of the nozzle, the design of heating chambers, the geometry/location of powder feed, and the cooling, either by water or air, are the major features which differ. The low residual coating stress produced in the continuous-combustion HVOF process allows significantly greater thickness capability than the plasma method, while providing lower porosity, lower oxide content, and higher coating adhesion. Coatings produced by the continuous-combustion HVOF process also have much better machinability compared with other methods, and coating porosity has closely approached wrought materials, as verified by gas permeability testing. Continuous-combustion HVOF systems are available with closed-loop computer control and robotics capability. Figures 20 and 21 compare the bond strengths and hardness of coatings deposited with various thermal spray processes. The correlation between hardness and particle velocity is shown in Fig. 22.

FIG. 20 BOND STRENGTHS OF TYPE 316 STAINLESS STEEL COATINGS DEPOSITED BY VARIOUS THERMAL SPRAY PROCESSES. JETKOTE IS A REGISTERED TRADENAME OF THERMODYNE, INC.; JP-5000 IS A REGISTERED TRADENAME OF HOBART TAFA TECHNOLOGIES, INC. SOURCE: REF 22

FIG. 21 HARDNESS OF WC-CO COATING VERSUS THERMAL SPRAY PROCESS. JP-5000 IS A REGISTERED TRADENAME OF HOBART TAFA TECHNOLOGIES, INC. SOURCE: REF 22

FIG. 22 HARDNESS OF WC-CO COATING VERSUS THERMAL SPRAY PROCESS PARTICLE VELOCITY. JP-5000 IS A REGISTERED TRADENAME OF HOBART TAFA TECHNOLOGIES, INC. SOURCE: REF 22

Surface Preparation (Ref 20) Thermal spray coatings rely primarily on mechanical bonding to the substrate surface. Because bonding relies on this mechanical interlocking of spray particles with the substrate surface, preparation of the surface is critical to a successful coating. Preparation of the substrate before spraying is virtually the same for all thermal spray processes. More detailed information on surface preparation can be found in Ref 23. Chemical Cleaning. Chemical cleaning is used on parts that are contaminated or impregnated with material that cannot

be removed by other methods. Surfaces can be chemically cleaned by vapor degreasing, steam, hot detergent washing, or application of an industrial cleaning solvent. Degreasing is usually the most economical and a fast way to remove lubricants and oils. Grit blasting is the most frequently used surface preparation procedure. The roughening cleans the surface of

contamination that may inhibit bonding and creates an irregular profile of minute surface irregularities. This enhances the adhesion of the coating. Before application of the thermal spray coating, the cleaned substrate should be blasted to a white-metal condition (according to Steel Structures Painting Council specification SSPC-SP5-82) with a minimum surface profile of 6 μm (250 μm) Ra (arithmetic average surface roughness). Sharp angular grit (18 to 24 mesh average) should be used. Typical-grit blasting materials include chilled iron grit, alumina (Al2O3), silica sand, crushed flint, crashed garnet, silicon carbide, and crushed slag. Rough threading is often used in conjunction with grit blasting for cylindrical surfaces. It consists essentially of

machining a thread form on the area to be sprayed. This type of preparation provides excellent adhesion of the coating to the substrate. Bond coats are often used when an extremely high bond strength is desired or when mechanical deformation of the

substrate is not possible or practical. This involves spraying materials that are exothermic and self-bonding. Nickelaluminum alloys (Ni-5 to 20Al) and molybdenum are the most commonly used bond materials. A thin coating (normally about 50 to 125 μm, or 2 to 5 mils) of bond material is generally applied before the thermally sprayed overcoat. Care should be taken to ensure that the bond material is compatible with the overcoat and that a galvanic cell will not be created, which may be detrimental to the performance of the coating system. Preheating of the substrate is often beneficial. Preheating to 65 to 95 °C (150 to 200 °F) will eliminate surface

condensation and will reduce shrink differentials between the coating and the substrate. It is generally employed only with the flame spray processes, which use their flame as the heat source for warming the substrate. Finishing Treatment

Thermal spray coatings usually exhibit two common features in the as-deposited condition, a sandpaper-like surface finish and a structure with inherent porosity. A typical surface roughness of an as-sprayed coating ranges from 5.0 to 13 m (200 to 500 μin.) Ra. The porosity usually ranges from 2 to 17 vol%, depending on the process by which the coating is deposited and the material sprayed. In many cases, thermal spray coatings are applied to machine elements where the coated component must conform to close dimensional tolerances or have a high surface finish. In many applications, coatings are exposed to corrosive solutions or hydraulic fluids, which can infiltrate the pores, resulting in fluid leakage or corrosion of the base material. These conditions can contribute to the premature failure of the coating. Many applications require the coating to be sealed before finishing by machining or grinding, to protect the thermal spray coating. Sealing is the process by which the pores of a coating are filled to eliminate the possibility of infiltration by fluids or

corrosive media that can contribute to premature failure. Sealing materials, which may be applied by brushing, spraying, or dipping, include petroleum-based waxes and organic materials such as phenolics, epoxy resins, polyesters, polyurethanes, and polyimides. Coating Finishing Method. Thermal spray coatings are finished to dimension either by single-point machining or by

grinding. Hard surfacing materials such as ceramics and carbides are restricted to grind finishing, typically with diamond wheels. Thermal spray coatings are sufficiently different from the same material in wrought form that different grinding wheel and finishing tool recommendations are almost always required. Therefore, tools and wheels should not be chosen based on experience with the parent material in wrought or cast form. Most thermal spray coatings exhibit some degree of porosity. However, the HVOF method produces coatings closest in machining and grinding behavior to wrought materials. In general, thermal spray coatings can be machined to 1 to 2 μm (40 to 80 μin.) Ra and ground to 0.25 to 0.5 μm (10 to 20 μin.) Ra. Many plasma spray ceramic coatings can be lapped down to 0.025 to 0.05 μm (1 to 2 μin.) Ra.


1. K.C. ANTONY ET AL., HARDFACING, WELDING, BRAZING, AND SOLDERING, 9TH ED., VOL 6, METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 771-793 9. K.M. KULKARNI AND V. ANAND, METAL POWDERS USED FOR HARDFACING, POWDER METALLURGY, 9TH ED., VOL 7, ASM HANDBOOK, AMERICAN SOCIETY FOR METALS, 1984, P 823-836 19. B.A. KUSHNER AND E.R. NOVINSKI, THERMAL SPRAY COATINGS, FRICTION, LUBRICATION, AND WEAR TECHNOLOGY, VOL 18, ASM HANDBOOK, ASM INTERNATIONAL, 1992, P 829-833 20. R.H. UNGER, THERMAL SPRAY COATINGS, CORROSION, 9TH ED., VOL 13, ASM HANDBOOK, ASM INTERNATIONAL, 1987, P 459-462 21. M.L. THORPE, THERMAL SPRAY: INDUSTRY IN TRANSITION, ADV. MATER. PROCESS., VOL 143 (NO. 5), 1993, P 50-56 22. G. IRONS, "HIGHER VELOCITY THERMAL SPRAY PROCESSES PRODUCE BETTER AIRCRAFT ENGINE COATINGS," SAE TECHNICAL PAPER 920947, PRESENTED AT THE 28TH ANNUAL AEROSPACE/AIRLINE PLATING AND METAL FINISHING FORUM AND EXPOSITION (SAN DIEGO, CA), APRIL 1992 23. F.N. LONGO, PROCESSING AND DESIGN, LESSON 3, COURSE 51, THERMAL SPRAY TECHNOLOGY, MATERIALS ENGINEERING INSTITUTE, ASM INTERNATIONAL, 1992 Hardfacing, Weld Cladding, and Dissimilar Metal Joining J.R. Davis, Davis & Associates

Thermal Spray Materials and Applications

The most outstanding feature of thermal spray coatings is probably their diverse applicability. There are two primary reasons for this. First, materials selection is almost unlimited. In fact, more than 200 coating materials with different characteristics of toughness, coefficient of friction, hardness, and other properties are available. These materials can be grouped as follows: • • • • • • •

PURE METALS METAL ALLOYS CERMETS CERAMICS CARBIDES POLYMERS SPECIAL COMPOSITE MATERIALS SUCH AS NICKEL-ALUMINUM AND NICKEL-GRAPHITE COMPOSITE POWDERS

Second, the thermal spray process, properly controlled, imparts very little heat to the substrate (100 to 260 °C, or 200 to 500 °F), avoiding metallurgical change, distortion, and oxidation. This allows thermal spray coatings to be applied to practically any substrate (metals, plastics, composites, etc.). Also, such coatings can often be applied to the finish machined part, a practice unthinkable with welding, heat treating, or other high-temperature processes. Thermal spray coatings provide the design engineer with many options. They offer material properties not available in wrought metals. Coatings can be applied to selected areas rather than treating the entire part. Manufacturing costs can be reduced by eliminating unnecessary processing steps. Extraordinary performance characteristics can be designed into a part to extend its useful life, creating a new, marketable product in the process. Thermal spray coatings provide the solution to many mechanical, electrical, and corrosion resistance problems involving metal parts and assemblies. However, there are certain applications where such coatings should not be used. Before a thermal spray coating is specified, its suitability can usually be determined according to these criteria (Ref 24): •

•

•

•

NO STRENGTH IS IMPARTED TO THE BASE MATERIAL BY A SPRAYED DEPOSIT. THE COMPONENT TO BE SPRAYED MUST, IN ITS PREPARED FORM, BE ABLE TO WITHSTAND ANY MECHANICAL LOADING THAT WILL BE EXPERIENCED IN SERVICE. (IN A FEW APPLICATIONS, SOME STRENGTH CAN BE ADDED BY A THERMAL SPRAY COATING; HOWEVER, SUCH USES ARE UNUSUAL AND SHOULD BE CAREFULLY TESTED.) IF THE AREA ON A PART TO BE SPRAYED OR ANY SECTION OF THE TOTAL AREA WILL BE SUBJECTED TO SHEAR LOADING IN SERVICE, THE PART IS NOT A SUITABLE CANDIDATE FOR THERMAL SPRAYING. GEAR TEETH, SPLINES, AND THREADS ARE EXAMPLES. POINT LOADING WITH LINE CONTACT ON A SPRAYED METAL DEPOSIT WILL EVENTUALLY SPREAD THE DEPOSIT, CAUSING DETACHMENT. IF THE DEPOSIT IS ON A MOVING COMPONENT WITH SUCH LOADING, IT WILL FAIL RAPIDLY. FOR EXAMPLE, NEEDLE AND ROLLER BEARING SEATS, WHERE THE BEARING ELEMENTS ARE IN DIRECT CONTACT WITH THE SPRAYED DEPOSIT, MAY NOT BE GOOD THERMAL SPRAY CANDIDATES. IF THE BASE METAL OF A COMPONENT TO BE TREATED HAS BEEN NITRIDED, NONHVOF THERMAL SPRAYING IS NOT RECOMMENDED UNLESS THE NITRIDING IS FIRST REMOVED. OTHER FORMS OF SUBSTRATE HARDENING REQUIRE SPECIAL TREATMENT, SUCH AS INTENSIVE GRIT BLASTING. THERMAL SPRAYING OF NITRIDED TURBINE ENGINE COMPONENTS USING HVOF METHODS IS DESCRIBED IN REF 25.

To date, the most common application area for thermal spraying has been advanced aircraft gas turbine components such as compressor blades, compressor stator vanes, beating housings, and labyrinth seals. Common coating materials include cobalt-base (Laves phase) alloys T-400 and T-800, tungsten carbide-cobalt materials, and chromium carbide-nickel chromium cermets (Cr3C2-20 to 30% NiCr). Other key application areas are listed below.

Wear coatings are used to resist abrasion, erosion, cavitation, and fretting, and to reduce friction. These coatings consist

of a wide range of metals and their alloys, ceramics, cermets, carbides, and even low-friction plastics. Typical coating hardness ranges from 20 to 70 HRC. Metal matrices are hardened by rapid solidification, by dispersions introduced from the spray process, by the addition of separate hard phases (for example, carbides of chromium, tungsten, titanium and/or tantalum), or by oxide inclusions. Table 13 lists friction and wear applications for various thermal spray materials.

TABLE 13 THERMAL SPRAY COATINGS USED FOR HARDFACING APPLICATIONS

TYPE OF WEAR ADHESIVE WEAR

COATING MATERIAL SOFT BEARING COATINGS: ALUMINUM BRONZE TOBIN BRONZE BABBITT TIN HARD BEARING COATINGS: MO/NI-CR-B-SI BLEND MOLYBDENUM HIGH-CARBON STEEL

CO-MO-CR-SI FE-MO-C ABRASIVE WEAR

SURFACE FATIGUE WEAR FRETTING:

ALUMINA/TITANIA TUNGSTEN CARBIDE PA, HVOF PA ALUMINUM OXIDE CHROMIUM OXIDE TUNGSTEN CARBIDE CHROMIUM CARBIDE NI-CR-B-SIC/WC (FUSED) NI-CR-B-SIC (FUSED) NI-CR-B-SIC (UNFUSED)

MOLYBDENUM

COATING PROCESS(A)

APPLICATIONS

OFW, EAW, OFP, PA, HVOF OFW, EAW OFW, EAW, OFP OFW, EAW, OFP

BABBITT BEARINGS, HYDRAULIC PRESS SLEEVES, THRUST BEARING SHOES, PISTON GUIDES, COMPRESSOR CROSSHEAD SLIPPERS

PA

BUMPER CRANKSHAFTS FOR PUNCH PRESS, SUGAR CANE GRINDING ROLL JOURNALS, ANTIGALLING SLEEVES, RUDDER BEARINGS, IMPELLER SHAFTS, PINION GEAR JOURNALS, PISTON RINGS (INTERNAL COMBUSTION); FUEL PUMP ROTORS

OFW, EAW, PA OFW, EAW OFP, PA OFP, PA, HVOF

PA PA PA, HVOF PA, HVOF OFP, HVOF OFP, HVOF

SLUSH-PUMP PISTON RODS, POLISH ROD LINERS, AND SUCKER ROD COUPLINGS (OIL INDUSTRY); CONCRETE MIXER SCREW CONVEYORS; GRINDING HAMMERS (TOBACCO INDUSTRY); CORE MANDRELS (DRY-CELL BATTERIES); BUFFING AND POLISHING FIXTURES; FUELROD MANDRELS

HVOF

OFW, PA

SERVOMOTOR SHAFTS, LATHE

COMBUSTION), CYLINDER LINERS FRETTING: SMALL-AMPLITUDE OSCILLATORY DISPLACEMENT APPLICATIONS: LOW TEMPERATURE (540 °C, OR 1000 °F)

CO-CR-NI-W CHROMIUM CARBIDE

PA, HVOF PA, HVOF

EROSION

CHROMIUM CARBIDE TUNGSTEN CARBIDE WC/NI-CR-B-SIC (FUSED) WC/NI-CR-B-SIC (UNFUSED) CHROMIUM OXIDE NI-CR-B-SIC-AL-MO NI-AL/NI-CR-B-SIC TYPE 316 STAINLESS STEEL NI-CR-B-SIC (FUSED) NI-CR-B-SIC (UNFUSED) ALUMINUM BRONZE CU-NI

PA, HVOF

CAVITATION

PA, HVOF OFP, HVOF

AIRCRAFT FLAP TRACKS (AIRFRAME COMPONENT); EXPANSION JOINTS AND MIDSPAN SUPPORTS (JET ENGINE COMPONENTS) COMPRESSOR AIR SEALS, COMPRESSOR STATORS, FAN DUCT SEGMENTS AND STIFFENERS (ALL JET ENGINE COMPONENTS) EXHAUST FANS, HYDROELECTRIC VALVES, CYCLONE DUST COLLECTOR, DUMP VALVE PLUGS AND SEATS, EXHAUST VALVE SEATS

OFP, HVOF PA PA PA PA OFP, HVOF

WEAR RINGS (HYDRAULIC TURBINES), WATER TURBINE BUCKETS, WATER TURBINE NOZZLES, DIESEL ENGINE CYLINDER LINES PUMPS

HVOF PA, HVOF PA, HVOF

(A) OFW, OXYFUEL WIRE SPRAY; EAW, ELECTRIC ARC WIRE SPRAY; OFP, OXYFUEL POWDER SPRAY; PA, PLASMA ARC SPRAY; HVOF, HIGH-VELOCITY OXYFUEL POWDER SPRAY Corrosion Resistance. Zinc and aluminum and their alloys are the metals most widely used for thermal spray

anticorrosion coatings. They are extensively used for the corrosion protection of iron and steel in a wide range of environments and have been shown to provide very long-term protection (over 20 years) in both marine and industrial locations. Properties and applications of zinc and aluminum thermal spray coatings are described in the article "Thermal Spraying Coatings" in this Volume. Other coating materials for corrosion-resistant applications include austenitic stainless steels, aluminum bronze, nickelbase alloys, and MCrAlY materials. Selection of a specific alloy depends on the particular environment. When using these alloys, it must be understood that these coatings will not galvanically protect the underlying steel. This can be a particular problem due to the porosity of thermal spray coatings. Great care must be taken to ensure that these coatings are properly sealed in order to prevent penetration of the corrosive medium to the underlying steel and subsequent corrosion at the coating/substrate interface.

Oxidation Protection. Thermal spray coatings are extensively used by industry to protect steel components and

structures from heat oxidation at surface temperatures to 1095 °C (2000 °F). By ensuring long-term protection, thermal spray coatings show real economic advantages during the service lives of such items. Coatings are particularly effective in protecting low-alloy and carbon steels. Table 14 lists thermal spray coatings for use in high-temperature applications.

TABLE 14 THERMAL SPRAY COATINGS USED FOR ELEVATED-TEMPERATURE SERVICE

SERVICE TEMPERATURE

COATING METAL OR ALLOY

UP TO 550 °C (1020 °F) UP TO 550 °C (1020 °F) IN THE PRESENCE OF SULFUROUS GASES 550-900 °C (1020-1650 °F)

ALUMINUM NI-43CR-2TI

900-1000 °C (1650-1830 °F) 900-1000 °C (1650-1830 °F) IN THE PRESENCE OF SULFUROUS GASES

ALUMINUM OR ALUMINUMIRON NICKEL-CHROMIUM OR MCRALY NICKEL-CHROMIUM, FOLLOWED BY ALUMINUM

COATING THICKNESS μm mils 175 7 375 15 175

7

375

15

375 100

15 4

Source: Ref 20

Aluminum has been widely used to protect such steels in a number of applications involving high surface temperatures up to 550 °C (1020 °F). Nickel-chromium alloys and some of the MCrAlY alloys have also provided protection in severe environments (up to 1000 °C, or 1830 °F). Thermal insulation provides low-thermal-conductivity zirconium and/or aluminum oxides that are deposited as coatings

to reduce heat conduction to base metals. These insulating oxide coatings can also reduce heat loss, reduce thermal transient effects, and reduce the high-temperature oxidation/corrosion effects by lowering temperatures of the coated base metals. Thermal spray processes usually introduce a controlled level of porosity, which further enhances the coating insulation and increases its thermal shock resistance. Thermal barrier coatings consist of a low-conductivity (thermal) ceramic deposited over an MCrAlY bond coat. The

ceramic of choice is partially-stabilized zirconia (7 to 8 wt% Y2O3-ZrO2) deposited at a thickness of 0.25 mm (0.010 to 0.040 in.) with 10 to 15% porosity. A bond coat of NiCrAlY or CoCrAlY is used at a thickness of 0.125 mm (0.005 in.). Both coatings are used on components used in gas turbine engines and adiabatic engines to improve efficiency and reduce metal temperatures or cooling requirements. Electrically conductive coatings include silver, copper, aluminum, zinc, and bronze alloys. Typical conductivities of

these sprayed metals are reported to be 40 to 90% of wrought materials, depending on the spray coating method. In addition, superconducting plasma-sprayed Y-Ba-Cu-O, Cu/Nb3Sn, and Nb3Sn materials that have zero resistance at low to intermediate temperatures of 100 K (-280 °F) have been deposited by thermal spraying. Conductive coatings designed to shield electronic components against radio frequency or electromagnetic interference such as aluminum, zinc, and tin coatings, are also sprayed. Dimensional Restorative Coatings. Thermal spray is used as a coating to repair or resurface (up to 3 mm, or 0.12 in.

thick) worn and/or corroded parts, as well as to repair new part errors (for example, undersize parts), to repair nicks or other blemishes, or to renew surfaces that are corroded or worn in use. Medically Compatible Coatings. Interactive bonding coatings on implants use thermal spray coatings to provide

biocompatible titanium (commercially pure Ti or Ti-6Al-4V alloy) and/or ceramic (calcium phosphates) surfaces on both dental and orthopedic devices, to improve wear surface and/or tissue adhesion. Coating porosity and high surface roughness can be produced in thermal spray coatings to improve further tissue adherence.


20. R.H. UNGER, THERMAL SPRAY COATINGS, CORROSION, 9TH ED., VOL 13, ASM HANDBOOK, ASM INTERNATIONAL, 1987, P 459-462 24. M.L. THORPE, THERMAL SPRAYING BECOMES A DESIGN TOOL, MACH. DES., 24 NOV 1983 25. P.O. BUZBY AND J. NIKITICH, HVOF THERMAL SPRAYING OF NITRIDED PARTS, ADV. MATER. PROCESS., DEC 1991, P 35-36 Hardfacing, Weld Cladding, and Dissimilar Metal Joining J.R. Davis, Davis & Associates

Weld Cladding The term weld cladding usually denotes the application of a relatively thick layer ([ges]3 mm, or

1 in.) of weld metal for 8

the purpose of providing a corrosion-resistant surface. Hardfacing, as described earlier in this article, produces a thinner surface coating than a weld cladding and is normally applied for dimensional restoration or wear resistance. Typical base metal components that are weld-cladded include the internal surfaces of carbon and low-alloy steel pressure vessels, paper digesters, urea reactors, tubesheets, nuclear reactor containment vessels, and hydrocrackers. The cladding material is usually an austenitic stainless steel or a nickel-base alloy, although certain copper-base alloys are also sometimes used. Some specialized cladding is also carried out using silver filler metal. Weld cladding is usually performed using the submerged arc welding process (SAW). However, flux-cored (either self-shielded or gas-shielded), plasma arc, and electroslag welding methods can also produce weld claddings. Filler metals are available as covered electrodes, coiled electrode wire, and strip electrodes. Table 15 lists some of the filler metals for stainless steel weld claddings.

TABLE 15 STAINLESS STEEL FILLER METALS FOR WELD CLADDING APPLICATIONS

WELD OVERLAY TYPE 304 304L 321 347 309 310 316 316L 317 317L 20 CB

FIRST LAYER COVERED ELECTRODE(A) E309 E309L E309CB E309CB E309CB E309 E310 E309MO E309MOL E317L E309MO E317 E309MOL E317L E320

BARE ROD OR ELECTRODE(B) ER309 ER309L ER309CB ER309CB ER309 ER310 ER309MO ER309MOL ER317L ER309MO ER317 ER309MOL ER317L ER320

Note: Columbium (Cb) is also referred to as niobium (Nb).

(A) (B)

REFER TO AWS SPECIFICATION A5.4. REFER TO AWS SPECIFICATION A5.9.

Application Considerations (Ref 26)

SUBSEQUENT LAYERS COVERED BARE ROD OR (A) ELECTRODE(B) ELECTRODE E308 ER308 E308L ER308L E347 E347 E309 E310 E316 E316L

ER347 ER347 ER309 ER310 ER316 ER316L

E317

ER317

E317L

ER317L

E320

ER320

The technique of weld cladding is an excellent method to impart properties to the surface of a substrate that are not available from that base metal, or to conserve expensive or difficult-to-obtain materials by using only a relatively thin surface layer on a less expensive or abundant base material. This technique has several inherent limitations or possible problems that must be considered when planning for weld cladding. The thickness of the required surface must be less than the maximum thickness of the overlay that can be obtained with the particular process and filler metal selected. Welding position also must be considered when selecting an overlay material and process. Certain processes are limited in their available welding positions; for example, SAW can be used only in the flat position. In addition, when using a high-deposition-rate process that exhibits a large liquid pool, welding vertically or overhead may be difficult or impossible. Some alloys exhibit eutectic solidification, which leads to large molten pools that solidify instantly rather than with a "mushy" (liquid plus solid) transition. Such materials are also difficult to weld except in the flat position. Composition Control of Stainless Steel Weld Overlays (Ref 26) The economics of stainless steel weld cladding are dependent on achieving the specific chemistry at the highest practical deposition rate in a minimum number of layers. The fabricator selects the filler wire and welding process, whereas the purchaser specifies the surface chemistry and thickness, along with the base metal. The most outstanding difference between welding a joint and depositing an overlay is in the area of dilution. The percentage of dilution equals the amount of base metal melted (x) divided by the sum of base metal melted and filler metal added (x + y), the quotient of which is multiplied by 100. %dilution =

x x100 x+ y

(EQ 2)

For stainless steel cladding, a fabricator must understand how the dilution of the filler metal with the base metal affects the composition and metallurgical balance, such as the proper ferrite level to minimize hot cracking, absence of martensite at the interface for bond integrity, and carbon at a low level to ensure corrosion resistance. The prediction of the microstructures and properties (such as hot cracking and corrosion resistance) for the austenitic stainless steels has been the topic of many studies (Ref 27). During the last two decades, four microstructure prediction diagrams have found the widest application. These include the Schaeffler diagram, the DeLong diagram, and the Welding Research Council (WRC) diagrams (WRC-1988 and WRC-1992). Each of these is described below. The Schaeffler diagram (Fig. 23), published in 1949, has been extensively used for estimating the ferrite content of

stainless steel weld metals (Ref 28). The diagram contains phase fields and iso-ferrite lines that permit prediction of the weld structure from the composition. It is still widely used for predicting the ferrite content of dissimilar weld deposits. This involves calculating a "chromium equivalent" (Creq) and a "nickel equivalent" (Nieq) for each base metal and for the proposed filler metal, plotting each equivalent on the Schaeffler diagram, and drawing tie lines between the plotted points, proportioned according to expected dilution, to obtain an estimate of the weld metal ferrite content.

FIG. 23 SCHAEFFLER DIAGRAM FOR STAINLESS STEEL WELD METAL. A, AUSTENITE; F, FERRITE; M, MARTENSITE. POINT 100 IS ASTM A 387A STEEL. EXAMPLE: POINT X ON THE DIAGRAM INDICATES THE EQUIVALENT COMPOSITION OF TYPE 318 (316 CB) WELD DEPOSIT CONTAINING 0.07% C, 1.55% MN, 0.57% SI, 19.02% CR, 11.87% NI, 2.16% MO, AND 0.80% NB. EACH OF THESE PERCENTAGES WAS MULTIPLIED BY THE "POTENCY FACTOR" INDICATED FOR THE ELEMENT IN QUESTION ALONG THE AXES OF THE DIAGRAM TO DETERMINE THE CHROMIUM AND THE NICKEL EQUIVALENT. WHEN THESE WERE PLOTTED AS POINT X, THE CONSTITUTION OF THE WELD WAS INDICATED AS AUSTENITE PLUS FROM 0 TO 5% FERRITE; MAGNETIC ANALYSIS OF THE ACTUAL SAMPLE REVEALED AN AVERAGE FERRITE CONTENT OF 2%. FOR AUSTENITE-PLUSFERRITE (A+F) STRUCTURES, THE DIAGRAM PREDICTS THE PERCENTAGE OF FERRITE WITHIN 4% FOR THE FOLLOWING STAINLESS STEELS: 308, 309, 309 CB, 310, 312, 316, 317, 318 (316 CB), AND 347. SOURCE: REF 26

As an example of the use of the Schaeffler diagram, suppose a type 304L clad surface is specified (shown as a triangle in Fig. 23). A desirable composition is approximately 20% Cr equivalent and 11% Ni equivalent. If an ASTM A 387A plate (point 100 on the diagram) is specified, all combinations of weld deposit lie on a line connecting point 100 with a proposed filler metal. As an example, note the line connecting point 100 with type 310 stainless steel. The diagram shows that weld-deposit compositions with up to 53% dilution (53% of the length of the line connecting the base metal and the filler metal) will be fully austenitic. Although sufficient alloy might be present, at least for a first layer, the absence of ferrite may make the deposit sensitive to fissuring. The same procedure can be used to estimate the usefulness of other filler metals. Good metallurgical practice dictates that the first layer contain no martensite, shown as M in the diagram. This limits the dilution. If a process can be found that has less than 14% dilution, then a type 308L could be deposited with absence of martensite. On the other hand, if a first layer is made using type 312 stainless, any dilution factor less than 40% prevents the formation of martensite. Choosing a filler metal for the first layer that falls in the A + F region with 3 to 15% ferrite and no martensite meets the metallurgical requirements of a sound deposit. Proper composition and thickness may be achieved in a single layer, but more often a second layer is required. Special consideration must be given to carbon levels for the austenitic stainless steels. Figure 24 predicts the carbon in the deposit once the base metal carbon and cladding process dilution factors have been established. To illustrate its use, suppose a base metal plate with 0.20% C and a welding process that results in 30% dilution are established. The diagram indicates that the deposit chemistry will contain approximately 0.07% C. If the specification requires 0.04% C maximum, a second layer would be required to achieve the necessary carbon level. Using the same diagram with the first layer

composition of 0.07% C, and assuming the same dilution factor, the diagram predicts that a second layer will contain approximately 0.03% C. This would meet the specification of 0.04% C maximum.

FIG. 24 PERCENTAGE OF CARBON IN THE DEPOSIT AFTER BASE METAL CARBON AND CLADDING DILUTION FACTORS HAVE BEEN ESTABLISHED. SOURCE: REF 26

Several problems have been identified when using the Schaeffler diagram for predicting the physical and mechanical properties of a weld. The Schaeffler diagram does not consider the effect of nitrogen in promoting austenite at the expense of ferrite, and it is incorrect in its treatment of manganese. Manganese does not promote the high-temperature formation of austenite at the expense of ferrite, as predicted by the diagram, although manganese does stabilize austenite in its low-temperature transformation to martensite. Also, the Schaeffler diagram makes its predictions in terms of percent ferrite. This has been found to be imprecise, and the magnetically based Ferrite Number (FN) unit was developed for the specification and determination of ferrite content. The basis for the FN scale is described in the American Welding Society standard AWS 4.2, "Standard Procedures for Calibrating Magnetic Instruments to Measure the Delta Ferrite Content of Austenitic and Duplex Austenitic Ferritic Steel." The DeLong Diagram. A second widely used prediction diagram, the DeLong diagram (Fig. 25) was published in 1974

and incorporated some improvements (Ref 29). It has an FN scale and includes a coefficient for nitrogen in the Nieq, though the range of compositions is more limited than for the Schaeffler diagram.

FIG. 25 THE DELONG CONSTITUTION DIAGRAM, REVISED JANUARY 1973 TO CONVERT IT TO THE WELDING RESEARCH COUNCIL'S FERRITE NUMBER (FN) SYSTEM FOR WELD METALS

The difference in the applicability of the Schaeffler and DeLong diagrams is substantial (Ref 30). The DeLong diagram is a finely tuned subset of the Schaeffler range, designed specifically for the 300-series stainless steel welds containing small amounts of ferrite. Its widely accepted use with these alloys is documented by its inclusion in the AWS filler metal specifications A5.4 and A5.9 and in the American Society of Mechanical Engineers (ASME) Boiler and Pressure Vessel Code. Prediction of FN for compositions that fall outside the range of the DeLong diagram (for example, duplex stainless steels) is the most quantitative use of the Schaeffler diagram; however, its extension to fully ferritic compositions and the inclusion of martensitic regions provides much qualitative information about the general effect of alloying elements and characteristic mechanical properties of alloy families. In effect, the DeLong diagram traded improved prediction accuracy for a more limited composition range, but a range that included all the common austenitic stainless steel weld metals. However, it has been found to incorrectly handle manganese (as does the Schaeffler diagram) and to overestimate the FN of more highly alloyed compositions, such as type 309 (Ref 30). It also has had limited application to dissimilar metal joints. The WRC-1988 diagram (Fig. 26) overcomes many of the problems associated with the Schaeffler and DeLong diagrams (Ref 30). It was developed with data measured by the most recent definition of the FN scale. It removed the erroneous manganese coefficient from the Nieq, and it eliminated the systematic overestimation of FN for highly alloyed weld metals. While it covers a much broader range of compositions than does the DeLong diagram, it has a narrower composition range than the Schaeffler diagram, because it extends only over the composition range of the commercial alloys with which it was developed.

FIG. 26 THE WRC-1988 DIAGRAM, INCLUDING SOLIDIFICATION-MODE BOUNDARIES. SOURCE: REF 30

In the interval since the WRC-1988 diagram first appeared, it has been subjected to considerable evaluation and review. Independent data from more than 200 weld metals have confirmed the improved prediction accuracy of the WRC-1988 over that of its predecessors (Ref 31). The WRC-1988 diagram predicts the ferrite content of the conventional 300-series stainless steels with a higher degree of accuracy than the DeLong diagram, is more accurate than the Schaeffler diagram for predicting the ferrite content of duplex stainless steels, and is suitable for new alloys containing extended compositional ranges. It also includes primary solidification-mode data to reveal relationships to the FN response. Recent studies have emphasized that the solidification mode is more accurate than the FN for predicting the resistance to hot cracking (Ref 30). These studies indicate that the initial solidification structure (primary austenite or ferrite) directly determines the sensitivity to cracking by controlling the solubility of elements that promote cracking at these temperatures. The WRC-1992 Diagram. In recent years, duplex ferritic-austenitic stainless steels have been used more frequently.

Some of these steels and their welds contain significant amounts of copper (often about 2 wt%). While these compositions might be plotted on the WRC-1988 diagram, greater accuracy can be obtained with the modified WRC-1992 diagram (Ref 32). Whereas the Schaeffler diagram offers predictions for Creq from 0 to 40 and Nieq from 0 to 32, the WRC-1992 diagram is more limited, considering Creq from 17 to 31 and Nieq from 9 to 18 (Fig. 27). As indicated in the following example, the WRC-1992 diagram is useful for predicting the FN for dissimilar metal joints.

FIG. 27 THE WRC-1992 DIAGRAM. SOURCE: REF 32

Example 1: Weld Cladding of AISl 1050 Steel with AWS A5.4 Class E312-16 Electrode (Ref 32). Table 16 lists typical plate composition and all-weld-metal composition produced with this electrode. The Creq and Nieq for the AISI 1050 steel (0.0 and 17.50, respectively) do not permit an FN calculation for this material. The Creq and Nieq for the E312-16 electrode all-weld-metal composition (29.0 and 11.9, respectively) result in a predicted 88.2 FN for this material. If the Creq and Nieq for each composition is plotted in Fig. 28 and a line drawn between the point corresponding to the AISI 1050 composition (Point A in Fig. 28) and the point corresponding to the E312-16 composition (Point B in Fig. 28), then all possible mixtures of these two materials must lie along this line.

TABLE 16 BASE METAL AND WELD METAL COMPOSITIONS PRODUCED WHEN AISI 1050 CARBON STEEL WAS WELD CLAD USING E312-16 STAINLESS STEEL FILLER METAL See Example 1 in text.

MATERIAL

AISI 1050 C, % 0.50 MN, % 0.30 SI, % 0.02 CR, % ... NI, % ... N, % 0.004 CREQ 0.00 NIEQ 17.58 WRC-1988 FN . . . Source: Ref 32

E312-16 ALL-WELD-METAL 0.060 1.20 0.60 29.00 8.60 0.06 29.00 11.90 88.2

30% DILUTION CLADDING 0.192 0.93 0.43 20.30 6.02 0.043 20.30 13.60 4.6

FIG. 28 DILUTION CALCULATION IN EXAMPLE 1 USING THE WRC-1992 DIAGRAM. SOURCE: REF 32

The lever rule is then used to estimate the position of the weld metal. In SMAW, typical dilution is 30%. A weld pass of E312-16 with 30% dilution from AISI 1050 steel would lie at a Point C, located 30% of the distance from Point B to Point A in Fig. 28. Mathematically, this can be calculated as:

CREQ(C) = 0.7 CREQ(B) + 0.3 CREQ(A) = 0.7(29.00) + 0.3(0.00) = 20.30

(EQ 3)

NIEQ(C) = 0.7 NIEQ(B) + 0.3 NIEQ(A) = 0.7(11.90) + 0.3(17.58) = 13.60

(EQ 4)

The Creq and Nieq of Point C correspond to 4.6 FN, which indicates that the weld pass will have sufficient FN to avoid fissuring (Ref 30). Any additional passes overlapping part of this first pass will also lie along the line between Point B and Point A, but they will be closer to Point B than is Point C, so all subsequent passes will contain even more ferrite than the first pass and the weld cladding should be crack-free. Control of dilution plays an important part in the economics of the weld cladding process. Although each process has an

expected dilution factor, experimenting with the welding parameters can minimize dilution. A value between 10 and 15% is generally considered optimum. Less than 10% raises the question of bond integrity, and greater than 15% increases the cost of the filler metal. Unfortunately, most welding processes have considerably greater dilution. Because of the importance of dilution in weld cladding as well as hardfacing applications, each welding parameter must be carefully evaluated and recorded. Many of the parameters that affect dilution in weld cladding applications are not so closely controlled when arc welding a joint. These parameters include: •

•

AMPERAGE: INCREASED AMPERAGE (CURRENT DENSITY) INCREASES DILUTION. THE ARC BECOMES HOTTER, IT PENETRATES MORE DEEPLY, AND MORE BASE METAL MELTING OCCURS. POLARITY: DIRECT CURRENT ELECTRODE NEGATIVE (DCEN) GIVES LESS PENETRATION AND RESULTING LOWER DILUTION THAN DIRECT CURRENT ELECTRODE POSITIVE

• •

•

•

•

•

•

(DCEP). ALTERNATING CURRENT RESULTS IN A DILUTION THAT LIES BETWEEN DCEN AND DCEP. ELECTRODE SIZE: THE SMALLER THE ELECTRODE, THE LOWER THE AMPERAGE, WHICH RESULTS IN LESS DILUTION. ELECTRODE EXTENSION: A LONG ELECTRODE EXTENSION FOR CONSUMABLE ELECTRODE PROCESSES DECREASES DILUTION. A SHORT ELECTRODE EXTENSION INCREASES DILUTION. TRAVEL SPEED: A DECREASE IN TRAVEL SPEED DECREASES THE AMOUNT OF BASE METAL MELTED AND INCREASES PROPORTIONALLY THE AMOUNT OF FILLER METAL MELTED, THUS DECREASING DILUTION. OSCILLATION: GREATER WIDTH OF ELECTRODE OSCILLATION REDUCES DILUTION. THE FREQUENCY OF OSCILLATION ALSO AFFECTS DILUTION--THE HIGHER THE FREQUENCY OF OSCILLATION, THE LOWER THE DILUTION. WELDING POSITION: DEPENDING ON THE WELDING POSITION OR WORK INCLINATION, GRAVITY CAUSES THE WELD POOL TO RUN AHEAD OF, REMAIN UNDER, OR RUN BEHIND THE ARC. IF THE WELD POOL STAYS AHEAD OF OR UNDER THE ARC, LESS BASE METAL PENETRATION AND RESULTING DILUTION WILL OCCUR. IF THE POOL IS TOO FAR AHEAD OF THE ARC, THERE WILL BE INSUFFICIENT MELTING OF THE SURFACE OF THE BASE METAL, AND COALESCENCE WILL NOT OCCUR. ARC SHIELDING: THE SHIELDING MEDIUM, GAS OR FLUX, ALSO AFFECTS DILUTION. THE FOLLOWING LIST RANKS VARIOUS SHIELDING MEDIUMS IN ORDER OF DECREASING DILUTION: GRANULAR FLUX WITHOUT ALLOY ADDITION (HIGHEST), HELIUM, CARBON DIOXIDE, ARGON, SELF-SHIELDED FCAW, AND GRANULAR FLUX WITH ALLOY ADDITION (LOWEST). ADDITIONAL FILLER METAL: EXTRA METAL (NOT INCLUDING THE ELECTRODE), ADDED TO THE WELD POOL AS POWDER, WIRE, STRIP, OR WITH FLUX, REDUCES DILUTION BY INCREASING THE TOTAL AMOUNT OF FILLER METAL AND REDUCING THE AMOUNT OF BASE METAL THAT IS MELTED.

Procedures for Stainless Steel Weld Cladding (Ref 26) Submerged arc welding is by far the most commonly used process for weld cladding. The process is adaptable for use

with single wires as the electrode filler metal, multiple wires, or strip. Alternating current is often used, but direct current, with either reverse or straight polarity, has been found preferable in numerous instances. Oscillation of the filler metal is frequently found desirable. Typical welding parameters for stainless steel weld cladding by SAW are given in Table 17.

TABLE 17 TYPICAL WELDING PARAMETERS FOR STAINLESS STEEL WELD CLADDING BY SAW FORM OF ELECTRODE

SIZE OF ELECTRODE

SINGLE WIRE

mm 1.6 2.4 2.4 3.2 4.0

SERIES ARC TWO WIRES

4.0

THREE WIRES

4.0

SIX WIRE FIRST WIRE

1.6

OTHERS

1.6

STRIP

0.5 × 30

POWER SOURCE

CURRENT, A

VOLTAGE, V

TRAVEL SPEED

in.

240

34

mm/min in./min mm 130 5 19

DCEN(B) 350

42

150

6

32

1

AC

350

30

280

11

AC

450

32

280

in. (B)

OSCILLATION

PRODUCTIVITY

DILUTION(A), %

kg/h 4

lb/h 9

9

20

15-20

NONE

5.0

11

20-25

11

NONE

5.4

12

20-25

1 16 3 32 3 32 1 8 5 32

DCEN

AC

500

34

280

11

NONE

5.9

13

20-25

5 32 5 32

AC(B)

400

26

255

10

NONE

19

41

20-25

AC(B)

480

28

355

14

NONE

20

43

15-20

1 16 1 16

DCEP

255

27

130

5

16

28

83

15-20

DCEN

1125

30

130

5

16

28

83

15-20

DCEP(B)

400

25

180

7

NONE

35

16

15-20

DCEP(B)

700

27

180

7

NONE

15

32

15-20

DCEP(B)

1250

27

130

5

NONE

27

60

15-20

DCEP(B)

1550

27

130

5

NONE

35

78

15-20

DCEP(B)

525

27

305

12

NONE

14

30

10-15

DCEP(B)

1050

27

305

12

NONE

27

60

10-15

0.020 × 1.2 0.5 × 61 0.020 × 2.4 0.5 × 88 0.020 × 3.5 0.5 × 0.020 × 119 4.7 0.64 × 0.025 × 1 25 0.64 × 0.025 × 2

3 4

5 8 5 8

1 4

15-20

50 0.64 100

× 0.025 × 4 DCEP(B)

2100

27

305

12

NONE

(A) VARIATIONS IN CURRENT, VOLTAGE, AND TRAVEL SPEED AFFECT PERCENTAGE OF DILUTION. (B) CONSTANT POTENTIAL

54

120

10-15

The single-wire process is largely used for overlaying limited areas or where it is necessary to limit the heat input. Wires from 1.6 to 4.8 mm (or

1 3 to in.) diam are practical. The small-size wires are generally employed with oscillation to 16 16

keep dilution low, to provide a wide (up to about 25 mm, or 1 in.) weld bead, and to give uniformly good tie-ins with the adjacent beads. Either alternating or direct current may be employed for single-wire overlays. Productivity (deposition rate) is greatly increased by the use of multiple wires (Table 17). Two-wire series arc (each wire connected to the poles of an alternating current power source) is a frequently used method for producing type 347 stainless steel weld cladding of petrochemical pressure vessels. The wire size is usually either 3.2 or 4.0 mm (

1 5 or in.) 8 32

in diameter. Alternating current is essential to achieve uniform fusion of both wires. The height of the welding head over the base metal greatly influences the penetration and hence the dilution. If the base metal has a non-uniform surface, or if the head cannot be accurately controlled above the base metal, the process may not be acceptable because of variations in dilution, or even lack of penetration in spots. "Shingling" is a technique used to minimize dilution in single- or two-wire weld claddings. By providing an overlap of as much as 70% of a bead width, it is possible to reduce the dilution from 30 or 35% to 15 or 20%. In some applications, this allows the cladding to be deposited in a single layer at considerable cost savings as compared with two-layer overlays. To further improve productivity, fabricators have increased the number of wires deposited by a single welding head. Up to six wires have been used with oscillation to deposit beads 150 mm (6 in.) or more in width. Usually, the wire sizes are 1.6 to 4.0 mm (

1 5 to in.) in diameter. Direct current electrode negative has been found best for relatively low dilution 16 32

(around 15%). Sometimes the wire that overlaps the previously deposited bead is connected to the positive tap of the power source (DCEP), which gives somewhat greater penetration and ensures good tie-in with the previous bead. For weld cladding the inside surfaces of large pressure vessels, as shown in Fig. 29(a)and 29(b), wide beads produced by oscillated multiple-wire systems or strip electrodes have become the means to improve productivity and minimize dilution while offering a uniformly smooth surface. Welding parameters for stainless steel strip weld overlays are illustrated in Fig. 30, 31, 32. In Fig. 30, the effect of increased current on dilution, penetration, and bead thickness of a 60 mm (2.4 in.) strip is shown. Figure 31 illustrates the effect of travel speed on dilution, penetration, and bead dimensions using a 60 mm (2.4 in.) strip. Productivity of different strip sizes (deposition rate versus arc current) is shown in Fig. 32.

FIG. 29(A) WELD CLADDING OF A 1.8 M (6 FT) INNER DIAMETER PRESSURE VESSEL SHELL WITH 50 MM (2 IN.) WIDE, 0.64 MM (0.025 IN.) THICK STAINLESS STEEL STRIP. COURTESY OF J.J. BARGER, ABB COMBUSTION ENGINEERING

FIG. 29(B) CLOSEUP VIEW OF THE 25 MM (1 IN.) WIDE BY 0.64 MM (0.025 IN.) THICK STAINLESS STEEL STRIP USED TO CLAD A 300 MM (12 IN.) INNER DIAMETER PRESSURE VESSEL NOZZLE. COURTESY OF J.J. BARGER, ABB COMBUSTION ENGINEERING

FIG. 30 DILUTION, PENETRATION, AND BEAD THICKNESS USING 60 MM (2.4 IN.) STRIP AT VARIOUS CURRENTS. SOURCE: REF 26

FIG. 31 INFLUENCE OF TRAVEL SPEED ON DILUTION, PENETRATION, AND BEAD THICKNESS USING 60 MM (2.4 IN.) STRIP. SOURCE: REF 26

FIG. 32 DEPOSITION RATE (100% DUTY CYCLE) AS A FUNCTION OF CURRENT AND DIFFERENT STRIP SIZES. SOURCE: REF 26

The fluxes for submerged arc overlays should have low bulk density, especially for wide weld beads produced with multiple wire or strip electrodes. Standard fused fluxes, therefore, are usually unsuitable. Agglomerated fluxes, which have much lower bulk density, are commonly used for this purpose. Low-bulk-density fused fluxes have been successfully produced by a manufacturing process known as foaming. Such foamed fluxes have even lower bulk density than agglomerated fluxes. Agglomerated fluxes (or bonded fluxes) have a distinct advantage over fused fluxes in that it is possible to introduce alloying elements by means of the flux. When this is done, they are known as reinforced, or alloyed, fluxes. For stainless steel weld claddings, chromium is commonly added to the fluxes to overcome the oxidation of chromium, which is characteristic of SAW of stainless alloys, and to compensate for the dilution with the non-chromium or low-chromium base metal. Other alloying elements often added to reinforced fluxes are manganese, nickel, molybdenum, and niobium. With reinforced fluxes, it is possible to produce various modified stainless weld overlays using a single filler metal as an electrode, for example, a type 347 overlay with a 308 or 309 filler metal using a niobium-alloyed flux. Table 18 lists variations in chemical compositions of series-arc weld deposits using an agglomerated flux and type 308 or 309 wire.

TABLE 18 CHEMICAL COMPOSITION VARIATIONS IN SERIES-ARC WELD CLADDING DEPOSITS USING AN AGGLOMERATED FLUX AND TYPE 308 OR 309 WIRE WELDING CONDITIONS

TYPE 308 OVERLAY TYPE 309 OVERLAY

WIRE SIZE mm in. 5 4.0 32 5 4.0 32

ANGLE BETWEEN WIRES, DEGREES

CURRENT (AC), A

VOLTAGE, V

TRAVEL SPEED mm/min in./min

45

500

20-25

255

10

45

540

26-30

330

13

CHEMICAL ANALYSIS COMPOSITION, % C Mn Si S 308 FILLER WIRE 0.04 1.74 0.30 0.011 TOP OR THIRD LAYER 0.05 1.39 0.61 0.007 SECOND LAYER 0.06 1.36 0.63 0.010 FIRST LAYER 0.07 1.20 0.58 0.006 CARBON STEEL BASE 0.07 0.51 . . . 0.013 309 FILLER WIRE 0.05 1.81 0.32 0.016 TOP OR THIRD LAYER 0.06 1.49 0.93 0.005 SECOND LAYER 0.06 1.55 0.84 0.006 FIRST LAYER 0.07 1.36 0.84 0.008 CARBON STEEL BASE 0.07 0.51 . . . 0.013

P 0.022 0.028 0.029 0.026 0.016 0.022 0.043 0.040 0.039 0.016

Cr 21.30 20.95 20.29 18.08 ... 24.68 24.43 24.37 21.79 ...

Ni 9.83 9.57 9.02 8.16 ... 13.36 13.14 11.08 11.43 ...

Another method for providing the desired stainless steel cladding composition is to introduce a layer of powdered metal under the submerged arc flux. This procedure is sometimes known as bulk welding. It is suitable for use either with single-wire or strip electrodes. The powdered metal not only provides the alloying elements needed in the overlay, but it also greatly reduces the dilution. A schematic of the process is shown in Fig. 33.

FIG. 33 SCHEMATIC OF THE BULKWELDING PROCESS USING METAL POWDER JOINT FILL. SOURCE: REF 9

Either method of adding alloying elements by using powdered alloys, either as an ingredient of the flux or separately underneath the flux, is subject to variability of the deposit composition. When reinforced fluxes are employed, the recovery of the alloying ingredient is dependent on the arc voltage, with a measurable difference in recovery by a change of as little as 1 or 2 V. This is not the case when the powdered alloys are separately fed below the flux, but in this case variability results if the height of the powder layer varies. Thus, it is unwise to rely on these methods of alloy adjustment of the weld overlay composition for more than three percentage points. Still another method of providing the desired composition for the cladding is to use tubular wire for the filler metal. Whereas the solid wire producer is limited in the compositions that can be produced and still provide hot workability and cold drawing capability, the producer of tubular wire is virtually unlimited as to the compositions that can be produced. For example, the strip can be chromium or chromium-nickel steel, to which the core ingredients are added to give a weld deposit of austenitic stainless steel, rich enough in the essential alloys to compensate for both oxidation and dilution losses. Another distinct advantage of tubular wires is for relatively small production applications of compositions that are not available as "shelf items" in solid-wire form. Small quantities of tubular wires can also be prepared for experiments to determine the optimum composition for a particular application. The effective cross section of tubular wire for currentcarrying ability is substantially less than solid wire of the same diameter. Welding parameters that have been established for solid wires will have to be adjusted to take this and other differences into account.

Self-Shielded Flux-Cored Wire. Tubular wires are produced that contain not only the alloying ingredients, but also the

fluxing and gas-shielding ingredients. This permits the deposition of weld claddings without the need of an external flux, as is needed with SAW, or an external gas, as is needed with gas-shielded solid wire welding. Weld claddings are produced by manually directed welding guns for limited areas using a single tubular filler wire. High productivity procedures have been developed using multiple wires in a manner similar to six-wire SAW. These wires have inherently low penetration characteristics, making them highly suitable for cladding. Deposition rates of 35 to 45 kg/h (80 to 100 lb/h) are often achieved. In addition, improved wire feed roll mechanisms have lessened the problems of kinking and crushing of these thin-wall electrodes during the wire-feeding process, thereby increasing the productivity of multiple-wire applications. Some fabricators who use self-shielded flux-cored wires prefer to limit the automatic cladding process to two or three wires, particularly because the gain in productivity with increasing numbers of wires diminishes. A three-wire welding head for overlay applications is shown in Fig. 34.

FIG. 34 THREE-WIRE WELDING HEAD FOR WELD CLADDING APPLICATIONS. COURTESY OF STOODY COMPANY

Metallurgically, flux-cored wires offer potentially higher alloyed claddings, permitting a considerable cost saving to produce the desired composition in a single layer. Also, dilution is minimized with this process. Furthermore, carbon content can be kept to much lower levels than is possible with solid wires, partly due to the ability to have very lowcarbon ingredients in the flux-cored wire and partly due to the low dilution. Typical welding parameters for stainless steel weld cladding using self-shielded flux-cored wires are given in Table 19.

TABLE 19 TYPICAL WELDING PARAMETERS FOR STAINLESS STEEL WELD CLADDING USING SELF-SHIELDED FLUX-CORED WIRES NO. OF 2.4 MM 8 ( IN.) DIAM 32 WIRES 1 2 3 6 (A)

CONSTANT POTENTIAL POWER CAPACITY(A)

CURRENT, A SOURCE

400 800 1200 1200 EACH (50 MM, OR 2 IN., PARALLEL) 100% DUTY CYCLE

VOLTAGE, V

300 600 900 1800

27 27 27 27

TRAVEL SPEED MM/MIN

OSCILLATION, CYCLES/MIN IN./MIN

510 115 100 90

20 4.5 4.0 3.5

NONE 20 20 20

PRODUCTIVITY KG/H

LB/H

5.4 14 20 39

12 30 45 85

DILUTION, %

20 12 12 12

The plasma arc hot wire process is used for depositing stainless steel claddings on pressure vessels subjected to

corrosive, high-temperature, and high-pressure hydrogen environments. The setup for this process is illustrated in Fig. 35. Using this method, two filler-metal wires are fed at a constant speed through individual hot wire nozzles to intersect in the weld pool beneath the plasma arc. The filler metals are electrically connected in series and energized by an alternating current constant-potential power supply. Resistance heating occurs in each electrode between the wire contact tip and the weld pool. Input power is selected for a given filler-metal feed rate so that maximum resistance heating is induced into both electrodes without resulting in an open arc between them. While a single hot wire increases the deposition rate considerably over the cold wire or powder methods, a dual hot wire process permits deposition rates as high as 18 to 32 kg/h (40 to 70 lb/h), thus making the hot wire process commercially competitive with other weld cladding processes. A continuously made stainless steel weld overlay made by the hot wire process is shown in Fig. 36.

FIG. 35 SCHEMATIC OF THE PLASMA ARC HOT WIRE WELD CLADDING PROCESS. SOURCE: REF 1

FIG. 36 PLASMA ARC HOT WIRE PROCESS FOR WELD CLADDING THE INTERNAL SURFACE OF A PRESSURE VESSEL SHELL. COURTESY OF NOOTER CORPORATION

Advantages of the hot wire process include: • • • •

ACCURATE CONTROL OF DILUTION AND THUS COMPOSITION OF THE OVERLAY LITTLE OR NO CHANGES DUE TO OXIDATION OF THE ELEMENTS IN THE PROCESS PARAMETERS THAT CAN BE INDEPENDENTLY AND ACCURATELY CONTROLLED SINGLE-LAYER OVERLAYS THAT CAN BE PRODUCED IN THICKNESSES RANGING FROM 4.0 TO 8.0 MM (

•

5 32

TO

5 IN.) 32

APPLICABILITY TO MANY CORROSION-RESISTANT ALLOYS, SUCH AS MONEL, INCONEL, AND HASTELLOY ALLOYS

Typical welding parameters for stainless steel weld cladding by the plasma arc process are given in Table 20.

TABLE 20 TYPICAL WELDING PARAMETERS FOR STAINLESS STEEL WELD CLADDING BY THE PLASMA ARC PROCESS TORCH CONDITIONS CURRENT VOLTAGE , , A V

GAS FLOW RATES mm3/ in.3/ h × h 106

FILLER-WIRE CONDITIONS SIZE NO. CURRENT OF , WIRE A m in. S m

440

38

1.4

85

2

1.6

480

38

1.4

85

2

1.6

500

39

1.4

85

2

1.6

500

39

1.4

85

2

2.4

1 16 1 16 1 16 3 32

TRAVEL SPEED

PRODUCTIVIT Y

DILUTION , %

mm/mi n

in./mi n

kg/h

lb/h

160

205

8

18-23

40-50

8-12

180

230

9

23-27

50-60

8-12

200

230

9

27-32

60-70

8-15

240

255

10

27-32

60-70

8-15

Electroslag Overlays. The electroslag welding process has been adapted to produce stainless steel weld claddings. This process offers the ability to use wider strip (up to 300 mm, or 12 in.) than the submerged arc process and more uniform penetration into, and lower dilution with, the base metal. Its principal disadvantages are operator exposure to the radiant heat from the molten slag and a sharper compositional change at the fusion line, which may give unsatisfactory service in certain high-temperature pressure vessels (for example, disbonding in hydrocracker or hydrosulfurizer service).

For large heavy-wall pressure vessels, the electroslag overlay process typically uses strip 4.0 mm (0.16 in.) thick by 150 mm (6 in.) wide. The effect of electroslag welding parameters on dilution using a 150 mm (6 in.) strip is shown in Fig. 37. As in SAW, energy is created by ohmic resistance as the electric current flows from the strip to the work through the molten slag; no arc exists. Unlike SAW, however, the slag bath is relatively thin. To provide the required electrical resistance, viscosity, and other properties, the fluxes must be significantly higher in fluorides than for SAW (Ref 26). A typical fused flux composition comprises 49% CaF2, 21% CaO, 21% SiO2, and 9% Al2O3.

FIG. 37 EFFECT OF WELDING PARAMETERS ON DILUTION USING 150 MM (6 IN.) STRIP IN ELECTROSLAG WELD CLADDING PROCESS. SOURCE: REF 26

Because of disbonding of weld claddings deposited on pressure vessels in a hydrogen environment, electroslag weld claddings are not used for single-layer overlays; but where two-layer overlayers are specified, the electroslag process is used for the second layer. Higher productivity and a smoother surface are thus obtained, the first layer having been deposited by the SAW. Extremely low carbon levels are possible because of low dilution with the base metal and a removal of carbon through oxidation in the process. Using a 0.18% C base metal and a 0.01% C strip, one investigator produced weld-metal carbon levels of 0.02% in the first layer and 0.01% in the second (Ref 26). The electroslag weld cladding process also can be applied for weld overlays of nickel-chromium (Inconel) alloys in the same manner as for stainless steel. Weld Overlays Other Than Stainless Steel (Ref 26) Many of the procedures described for stainless steel weld cladding are applicable to other alloys. Generally, nickel-base alloys can be produced successfully using essentially the same procedures as for stainless steels. Some of the common alloy filler metals used are given in Table 21. Figure 38 shows the use of both Inconel and stainless steel strip overlays used to clad a large pressure vessel shell.

TABLE 21 NONFERROUS FILLER METALS USED FOR WELD CLADDING OF STEELS WELD OVERLAY TYPE

ALUMINUM-BRONZE SILICON-BRONZE COPPER-NICKEL NICKEL MONEL LNCONEL LNCONEL 625 HASTELLOY B

COVERED ELECTRODE AWS CLASSIFICATION SPECIFICATION NO. NO. A5.6 ECUAL-A2 A5.6 ECUSI A5.6 ECUNI(A) A5.11 ENI-1 A5.11 ENICU-7 A5.11 ENICRFE-3 A5.11 ENICRMO-3 A5.11 ENIMO-7

BARE ROD OR ELECTRODE CLASSIFICATION AWS SPECIFICATION NO. NO. A5.7 ERCUAL-AL A5.7 ERCUSI A5.7 ERCUNI(A) A5.14 ERNI-1 A5.14 ENICU-7 A5.14 ERNICR-3 A5.14 ERNICRMO-3 A5.14 ERNIMO-7

HASTELLOY C

(A)

A5.11

ENICRMO-4

A5.14

ERNICRMO-4

FIRST LAYER MUST BE EITHER NICKEL OR MONEL.

FIG. 38 WELD CLADDING OF A 1.8 M (6 FT) INNER DIAMETER PRESSURE VESSEL SHELL USING 100 MM (4 IN.) WIDE BY 0.64 MM (0.025 IN.) THICK STAINLESS STEEL STRIP. NOTE SMALL "WINDOWS" THAT HAVE BEEN CLAD WITH 50 MM (2 IN.) WIDE BY 0.64 MM (0.025 IN.) THICK INCONEL STRIP. COURTESY OF J.J. BARGER, COMBUSTION ENGINEERING

Metals and alloys that do not tolerate significant amounts of iron cannot be used in weld overlays on steels. Titanium, zirconium, magnesium, and aluminum are examples of incompatible metals for cladding steel. Pure copper is also unsatisfactory, but some copper alloys, such as aluminum-bronze and silicon-bronze, can be applied to steels. The incompatible metals are usually bonded by explosion cladding or roll bonding techniques, which are described elsewhere in this Volume.


1. K.C. ANTONY ET AL., HARDFACING, WELDING, BRAZING, AND SOLDERING, 9TH ED., VOL 6, METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 771-793 9. K.M. KULKARNI AND V. ANAND, METAL POWDERS USED FOR HARDFACING, POWDER METALLURGY, 9TH ED., VOL 7, ASM HANDBOOK, AMERICAN SOCIETY FOR METALS, 1984, P 823-836 26. R.D. THOMAS ET AL., WELD OVERLAYS, WELDING, BRAZING, AND SOLDERING, 9TH ED., VOL 6, METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 804-819 27. D.L. OLSON, PREDICTION OF AUSTENITIC WELD METAL MICROSTRUCTURE AND PROPERTIES, WELD. J., VOL 64 (NO. 10), 1985, P 2815-2955 28. A.L. SCHAEFFLER, CONSTITUTION DIAGRAM FOR STAINLESS STEEL WELD METAL, MET. PROG., VOL 56 (NO. 11), 1949, P 680-680B 29. W.T. DELONG, FERRITE IN AUSTENITIC STAINLESS STEEL WELD METAL, WELD. J., VOL 53 (NO. 7), 1974, P 2735-2865 30. T.A. SIEWERT, C.N. MCCOWAN, AND D.L. OLSON, FERRITE NUMBER PREDICTION TO 100 FN IN STAINLESS STEEL WELD METAL, WELD. J., VOL 67 (NO. 12), 1988, P 2895-2985 31. D.J. KOTECKI, "VERIFICATION OF THE NBS-CSM FERRITE DIAGRAM," INTERNATIONAL INSTITUTE OF WELDING DOCUMENT II-C-834-88, 1988 32. D.J. KOTECKI AND T.A. SIEWERT, WRC-1992 CONSTITUTION DIAGRAM FOR STAINLESS STEEL WELD METALS: A MODIFICATION OF THE WRC-1988 DIAGRAM, WELD. J., MAY 1992, P 171S-178S Hardfacing, Weld Cladding, and Dissimilar Metal Joining J.R. Davis, Davis & Associates

Dissimilar Metal Joining Because various portions of a process system operate at different service conditions, different structural alloys are used in the design, and hence, dissimilar-metal welded joints may be required. Many factors must be considered when welding dissimilar metals, and the development and qualification of adequate procedures for the various metals and sizes of interest for a specific application must be undertaken. Most combinations of dissimilar metals can be joined by solid-state welding (diffusion welding, explosion welding, friction welding, or ultrasonic welding), brazing, or soldering where alloying between the metals is normally insignificant. In these cases, only the differences in the physical and mechanical properties of the base metals and their influence on the serviceability of the joint should be considered. When dissimilar metals are joined by fusion welding processes, alloying between the base metals and a filler metal, when used, becomes a major consideration. The resulting weld metal can behave much differently from one or both base metals during subsequent processing or in service. In this section, the principal factors that are responsible for failure (cracking) of dissimilar metal arc welds are described. These include: • • • •

GENERAL ALLOYING PROBLEMS (BRITTLE PHASE FORMATION AND LIMITED MUTUAL SOLUBILITY) OF THE TWO METALS WIDELY DIFFERING MELTING POINTS DIFFERENCES IN COEFFICIENTS OF THERMAL EXPANSION DIFFERENCES IN THERMAL CONDUCTIVITY

In addition, service considerations that can result in dissimilar-metal weld failure by carbon migration or by corrosion/oxidation are addressed. Lastly, a brief review of specific dissimilar metal combinations and filler metal selection guidelines are presented. Additional information on welding dissimilar metals can be found throughout this Volume and in Ref 33 and 34. Factors Influencing Joint Integrity (Ref 33) Weld Metal. In the fusion welding of dissimilar-metal joints, the most important consideration is the weld metal

composition and its properties. Its composition depends upon the compositions of the base metals; the filler metal, if used; and the relative dilutions of these. The weld metal composition is usually not uniform, particularly with multipass welds, and a composition gradient is likely to exist in the weld metal adjacent to each base metal. These solidification characteristics of the weld metal are also influenced by the relative dilutions and the composition gradients near each base metal. These characteristics are important with respect to hot cracking of the weld metal during solidification. The basic concepts of alloying, the metallurgical characteristics of the resultant alloy, and its mechanical and physical properties must be considered when designing a dissimilar-metal joint. For the fusion welding processes it is important to investigate the phase diagram of the two metals involved. If there is mutual solubility of the two metals, the joint can usually be made successfully. If there is little or no solubility between the two metals to be joined, the weld joint will not be successful. The intermetallic compounds that are formed between the dissimilar metals must be investigated to determine their crack sensitivity, ductility, susceptibility to corrosion, and so on. The microstructure of this intermetallic compound is extremely important. In some cases, it is necessary to use a third metal that is soluble with each metal in order to produce a successful joint. Dilution. In dissimilar-metal welding, the filler metal must alloy readily with the base metals to produce a weld metal

that has a continuous, ductile matrix phase. Specifically, the filler metal must be able to accept dilution (alloying) by the base metals without producing a crack-sensitive microstructure (dilution was defined and described earlier in this article in the sections on "Hardfacing" and "Weld Cladding" ). The weld metal microstructure must also be stable under the expected service conditions. A successful weld between dissimilar metals is one that is as strong as the weaker of the two metals being joined, that is, possessing sufficient tensile strength and ductility so that the joint will not fail.

In multipass welding, the composition of each weld bead should be relatively uniform. However, definite compositional differences are likely in succeeding weld beads, especially between a root bead, the beads adjacent to the base metals, and the remaining fill beads. The average composition of the whole weld metal can be calculated when: (1) the ratio of the volumes of base metals melted to the entire weld metal volume can be determined; and (2) the compositions of the base and filler metals are known. The dilution can be based on area measurements on a transverse cross section through a test weld. Figure 39 illustrates how to determine the dilution by two base metals, A and B, when welding with filler metal F.

FIG. 39 DILUTION IN A DISSIMILAR-METAL WELDED JOINT. SOURCE: REF 33

The average percentage of a specific alloying element in the diluted weld metal can be calculated using the following equation developed by the American Welding Society (Ref 33):

XW = (DA)(XA) + (DB)(XB) + (1 - DT)(XF)

(EQ 5)

where Xw is the average percentage of element X in the weld metal; XA is the percentage of element X in base metal A; XB is the percentage of element X in base metal B; XF is the percentage of element X in the filler metal F; DA is the percent dilution by base metal A, expressed as a decimal; DB is the percent dilution by base metal B, expressed as a decimal; and DT is the percent total dilution by base metals A and B, expressed as a decimal. To illustrate the calculation of weld metal composition, assume that type 316 stainless steel is welded to a 2.25Cr-1Mo (UNS K21590) low-alloy ferritic pressure vessel steel with a nickel-chromium alloy filler metal ERNiCr3. The nominal chemical compositions of the three alloys are:

COMPOSITION, WT% CR NI MO FE 2.25CR-1MO 2.5 . . . 1.0 95.5 TYPE 316 17.0 12.0 2.5 63.0 ERNICR-3 20.0 72.0 . . . 3.0 ALLOY

Assuming that the total dilution is 35%, 15% by the Cr-Mo alloy steel and 20% from the type 316 stainless steel, the average percentages of Cr, Ni, and Mo in the weld metal are calculated as follows:

CR, % = 0.15(2.5) + 0.20(17) + 0.65(20) = 16.8

(EQ 6)

NI, % = 0.20(12) + 0.65(72) = 49.2

(EQ 7)

MO, % = 0.15(1) + 0.20(2.5) = 0.65

(EQ 8)

Melting Temperatures. The difference in melting temperatures of the two metals that are to be joined must also be

considered. This is of primary interest when a fusion welding process utilizing considerable heat is involved, since one metal may be molten long before the other when subjected to the same heat source. Significant difference between the melting temperatures of the two base metals or between those of the weld metal and a base metal can result in rupture of the metal having the lower melting temperature. Solidification and contraction of the metal with the higher melting temperature will induce stresses in the other metal while it is in a weak, partially solidified condition. This problem may be solved by depositing one or more layers of a filler metal of intermediate melting temperature on the face of the base metal with the higher melting temperature. This procedure is known as buttering. The weld is then made between the buttered face and the other base metal. The buttering layer should serve to reduce the melting temperature differential. Buttering may also be used to provide a transition between materials with substantially different coefficients of thermal expansion but which must endure cycling temperatures in service. Similarly, buttering may be used to provide a barrier layer that will slow the migration of undesirable elements from the base metal to the weld metal during postweld heat treatment or in service at elevated temperatures. Thermal Conductivity. Most metals and alloys are relatively good conductors of heat, but some are much better than others. Rapid conduction of heat from the molten weld pool by an adjacent base metal may affect the energy input required to locally melt the base metal. When two dissimilar metals of significantly different thermal conductivities are welded together (for example, plain carbon steels and copper-base alloys), the welding procedure must provide for this difference. Often the welding heat source must be directed at the metal having the higher thermal conductivity to obtain the proper heat balance.

When welding dissimilar metals, heat loss to the base metals can be balanced somewhat by selectively preheating the metal having the higher thermal conductivity. Dilution is more uniform with balanced heating. Preheating the base metal of higher thermal conductivity also reduces the cooling rate of the weld metal and the HAZ. The net effect of preheating is to reduce the heat needed to melt that base metal. The coefficient of thermal expansion of the two dissimilar base metals is another important factor. Large differences

in thermal expansion coefficients of adjacent metals during cooling will induce tensile stress in one metal and compressive stress in the other. The metal subject to tensile stress may hot crack during welding, or it may cold crack in service unless the stresses are relieved thermally or mechanically. This factor is particularly important in joints that will operate at elevated temperatures in a cyclic temperature mode. A common example of this is austenitic stainless steel/ferritic steel pipe butt joints used in energy conversion plants. Ideally, the thermal expansion coefficient of the weld metal should be intermediate between those of the base metals, especially if the difference between those of the two base metals is large. If the difference is small, the weld metal may have an expansion coefficient equivalent to that of one of the base metals. Welding Considerations (Ref 33) Welding Process. The three most popular arc welding processes utilized for joining dissimilar metals are SMAW,

GMAW, and GTAW. Selecting the welding process to make a given dissimilar-metal joint is almost as important as selecting the proper filler metal. As described in the articles in this Volume that cover the various fusion welding processes, the depth of fusion into the base metals and the resulting dilution may vary with different welding processes and techniques. It is not uncommon with SMAW for the filler metal to be diluted up to 30% with base metal. The amount of dilution can be modified somewhat by adjusting the welding technique. For example, the electrode can be manipulated so that the arc impinges primarily on the previously deposited weld metal. The dilution rate can be kept below 25% with this technique. If dilution from one base metal is less detrimental than from the other, the arc should be directed toward that metal. This

technique is also applicable to GTAW. Other means to control dilution are described in the section on "Weld Cladding" in this article. Dilution rates with GMAW can range from 10 to 50%, depending upon the type of metal transfer and the welding gun manipulation. Spray transfer gives the greatest dilution; short-circuiting transfer, the least dilution. Penetration with SAW can be greater, depending on polarity, and can result in more dilution. Regardless of the process, dilution is also affected by other factors, including joint design and fit-up. It is always best to have a minimum uniform dilution along the joint. Variations in dilution may produce inconsistent joint properties. Selection of a suitable filler metal is an important factor in producing a dissimilar-metal joint that will perform well in service. One objective of dissimilar-metal welding is to minimize undesirable metallurgical interactions between the metals. The filler metal should be compatible with both base metals and be capable of being deposited with a minimum of dilution.

Two important criteria that should govern the selection of a proper filler metal for welding two dissimilar metals are as follows: • •

THE CANDIDATE FILLER METAL MUST PROVIDE THE JOINT DESIGN REQUIREMENTS, SUCH AS MECHANICAL PROPERTIES OR CORROSION RESISTANCE. THE CANDIDATE FILLER METAL MUST FULFILL THE WELDABILITY CRITERIA WITH RESPECT TO DILUTION, MELTING TEMPERATURE, AND OTHER PHYSICAL PROPERTY REQUIREMENTS OF THE WELDMENT.

The Schaeffler diagram is commonly used to predict weld metal microstructure and subsequent filler metal selection when joining a stainless steel to a carbon or low-alloy steel. Figure 40 illustrates the procedure with an example of a single-pass weld joining mild steel to type 304 stainless steel with ER309 stainless steel filler metal. First a connecting line is drawn between the two points representing the base metal compositions, based on their chromium and nickel equivalents. Point X, representing the relative dilutions contributed by each base metal is then located on this line. If the relative dilutions are equal, point X is at the midpoint of the line. A second line is drawn between point X and the point representing the ER309 filler metal composition. The composition of the weld pass lies somewhere on this line, the exact location depending upon the total dilution. With 30% dilution, the composition would be at point Y and would be considered acceptable. If a succeeding pass joins the first pass to mild steel, the dilution with the mild steel should be kept to a minimum, to avoid martensite formation in the weld metal.

FIG. 40 PREDICTION OF WELD METAL COMPOSITION FROM THE SCHAEFFLER DIAGRAM. A, AUSTENITE; F, FERRITE, M, MARTENSITE. SEE TEXT FOR DETAILS. SOURCE: REF 33

For improved accuracy in predicting the microstructure of dissimilar-metal welds, the WRC-1988 or WRC-1992 diagrams should be used. The following example demonstrates the use of the WRC-1992 diagram for predicting the microstructure of a structural steel/stainless steel weld. Example 2: Joining Type 304 Stainless Steel to ASTM A 36 Structural Steel with E309L16 Electrode (Ref 32). Table 22 lists typical plate compositions and all-weld-metal composition from the electrode. The Creq and Nieq for the A 36 steel (0.00 and 7.08, respectively) do not permit an FN calculation for this material. The Creq and Nieq for the type 304 stainless steel (18.83 and 12.45, respectively) result in a calculated 3.2 FN for this material. The Creq and Nieq for these two base metals are shown in Fig. 41 and Point D (A 36) and Point E (type 304). Any mixture of these two base metals will lie along the line connecting Point D to E. If each base metal contributes equally to the weld metal, then the overall base metal contribution is given by the midpoint of the line between Point D and Point E, indicated in Fig. 41 as Point F. If one base metal contributed more than the other to the joint (for example, due to unequal plate thickness of a complex joint design), the Point F would slide along this line proportionately toward the greater contributor. In any case, the average base metal contribution to the weld pool would be along this line.

TABLE 22 BASE METAL AND WELD METAL COMPOSITIONS PRODUCED WHEN ASTM A 36 STRUCTURAL STEEL AND TYPE 304 STAINLESS STEEL WERE JOINED BY SMAW USING E309L-16 FILLER METAL See Example 2 in text.

MATERIAL C, % MN, % SI, % CR, %

AISI 304 0.05 1.60 0.40 18.75

ASTM A 36 0.20 0.80 0.20 ...

E309L-16 ALL-WELD-METAL 0.03 1.40 0.60 24.40

70% E309L-16, 15% AISI 304, 15% A 36 0.059 1.34 0.51 19.89

NI, % MO, % N, % CREQ NIEQ WRC-1988 FN

9.90 0.08 0.04 18.83 12.45 3.2

... ... 0.004 0.00 7.08 ...

12.70 0.20 0.06 24.60 14.95 17.4

10.38 0.15 0.049 20.04 13.39 4.3

Source: Ref 32

FIG. 41 DILUTION CALCULATION IN EXAMPLE 2 USING THE WRC-1992 DIAGRAM. SOURCE: REF 32

The all-weld-metal Creq and Nieq (24.60 and 14.95, respectively) for the E309L-16 electrode is shown as Point G in Fig. 41, and a 17.4 FN would be predicted for that electrode. The root pass weld metal, consisting of the E309L-16 electrode and equal parts of the two base metals, must lie along the line from Point G to Point F in Fig. 41. Assuming the normal 30% base metal dilution with the shielded metal arc process, the root pass weld would lie 30% of the distance along the line from Point G to Point F. This is shown as Point H in Fig. 41. The calculations to reach Point H are as follows:

CREQ(F) = 0.5 CREQ(D) + 0.5 CREQ(E) = 0.5(0.00) + 0.5(18.83) = 9.47 NIEQ(F) = 0.5 NIEQ(D) + 0.5 NIEQ(E) = 0.5(7.08) + 0.5(12.45) = 9.77

(EQ 9) (EQ 10)

CREQ(H) = 0.7 CREQ(G) + 0.3 CREQ(F) = 0.7(24.60) + 0.3(9.47) = 20.04

(EQ 11)

NIEQ(H) = 0.7 NIEQ(G) + 0.3 NIEQ(F) = 0.7(14.95) + 0.3(9.77) = 13.39

(EQ 12)

Point H corresponds to 4.3 FN, which would be expected to result in a crack-free root pass in this joint. Higher dilution, or excess dilution from the A 36 side of the joint, could reduce the ferrite content and increase the likelihood of cracking.

Buttering. If dilution of austenitic stainless steel filler metal is a problem, it may be controlled by first buttering the joint

face of the carbon or low-alloy steel with one or two layers of type 309 or 310 stainless steel filler metal, as shown in Fig. 42(a) and 42(b). After machining and inspecting the buttered layer (Fig. 42c), the joint between the stainless steel component and the buttered steel part can be made using conventional welding procedures and the appropriate filler metal for welding the stainless steel base metal (Fig. 42d and 42e). A low-alloy steel component can be heat-treated after the buttering operation, then joined to the stainless steel part. This avoids a postweld heat treatment that might sensitize the austenitic stainless steel to intergranular corrosion. Additional information on buttering can be found in the article "Dissimilar Welds with Stainless Steels" in this Volume.

FIG. 42 BUTTERING TECHNIQUE USED TO ASSIST WELDING STAINLESS STEEL TO CARBON OR LOW-ALLOY STEEL. (A) EDGE PREPARED FOR BUTTERING. (B) FACE BUTTERED WITH FILLER METAL. (C) BUTTERED FACE PREPARED FOR WELDING. (D) JOINT ALIGNED FOR WELDING. (E) JOINT WELDED WITH STAINLESS STEEL FILLER METAL. SOURCE: REF 33

Joint Design. When designing butt joints between dissimilar metals, consideration must be given to the melting

characteristics of each base metal and the filler metal, as well as to dilution effect. Large grooves decrease dilution, permit better control of viscous weld metal, and provide room for better manipulation of the arc for good fusion. The joint design should provide for appropriate dilution in the first few passes placed in the joint when welding from one side. Improper dilution could result in a layer of weld metal possessing mechanical properties that are inappropriate for the intended service, particularly when the joint will be exposed to cyclic stresses. When welding from both sides, back gouging of the first weld can provide better control of dilution in the first few passes of the second weld. Dissimilar-metal joints are often made with groove geometries similar to those used in conventional similar-metal welds and are left in the as-welded condition. However, removing the excess penetration of the weld root eliminates possible defects that may be present in the root of the joint, and it eliminates notches and crevices associated with backing rings. Removing the weld root also improves the capabilities of nondestructive tests (radiography, liquid penetrant, and ultrasonic). Provision can also be made for removing the exterior reinforcement to eliminate stress concentrations associated with undercutting and with the change in section thickness across the as-welded joint. Again, inspection capabilities are improved. Preheat and Postweld Heat Treatments. Selection of an appropriate preheat or postweld heat treatment (PWHT) for a welded joint can present a problem with some dissimilar metal combinations. The appropriate heat treatment for one component of the weldment may be deleterious to the other component for the intended service conditions. For example,

if an age-hardenable nickel-chromium alloy is welded to a nonstabilized austenitic stainless steel, exposure of the weldment to the aging treatment for the nickel-chromium alloy would sensitize the stainless steel and decrease its resistance to intergranular corrosion (weld decay). One solution is to use a stabilized austenitic stainless steel if that is acceptable. Another solution might be to butter the face of the age-hardenable, nickel-chromium alloy component with a similar alloy that is not age-hardenable. This component is then heat-treated to obtain the desired properties. Finally, the buttered surface is welded to the stainless steel component. Service Considerations (Ref 33) Property Considerations. A dissimilar metal joint normally contains weld metal having a composition different from

that of one or both base metals. The properties of the weld metal depend on the filler metal composition, the welding procedures, and the relative dilution with each base metal. There also two different HAZs, one in each base metal adjacent to the weld metal. The mechanical and physical properties of the weld metal as well as those of the two HAZs, must be considered for the intended service. Special considerations are generally given to dissimilar-metal joints intended for elevated-temperature service. A favorable situation exists when the joint will operate at constant temperatures. During elevated-temperature service, internal stresses can decrease by relaxation and reach an equilibrium. However, it is best to reduce the effect of large differences in coefficients of thermal expansion when large temperature fluctuations cannot be avoided in service. The problem can be avoided by selecting base metals with similar thermal expansion characteristics. Figure 43 gives the mean linear thermal expansion coefficients as a function of temperature for several alloys commonly associated with transition butt joints for steam plant applications. The thermal expansion coefficient of 2.25Cr-1Mo steel is about 25% less than that of types 304 and 316 austenitic stainless steel. In certain applications, transition joints between austenitic stainless steel and low-alloy steel may experience numerous temperature changes during operation. For a given change in temperature, the stress imposed at the weld joint is proportional to the difference in their coefficients of thermal expansion. Stress analyses of welded joints between these two types of steel indicate that the stresses introduced by thermal change are nearly an order of magnitude greater than those produced by the operating pressures and any thermal gradients across the joint.

FIG. 43 MEAN COEFFICIENTS OF THERMAL EXPANSION AS A FUNCTION OF TEMPERATURE FOR TRANSITION BUTT-JOINT MATERIALS. SOURCE: REF 33

Experience with dissimilar-metal transition joints has shown a significant number of failures in less than the expected service life. The majority of the transition joint failures in austenitic steel/ferritic steel joints occur in the ferritic steel HAZ, adjacent to the weld interface. These failures are generally attributed to one or more of the following causes: •

• •

HIGH STRESSES AND THE RESULTING CREEP AT THE INTERFACE AS A RESULT OF THE DIFFERENCES BETWEEN THE COEFFICIENTS OF THERMAL EXPANSION OF THE WELD AND BASE METALS CARBON MIGRATION FROM THE FERRITIC STEEL INTO THE STAINLESS STEEL, WHICH WEAKENS THE HAZ IN THE FERRITIC STEEL PREFERENTIAL OXIDATION AT THE INTERFACE, ACCELERATED BY THE PRESENCE OF STRESS

Carbon Migration. There is the likelihood of significant chemical composition gradients in the weld metal, particularly

in the regions adjacent to the base metals. In addition, operation at elevated temperatures can cause interdiffusion between the weld metal and the base metal that, in turn, results in microstructure changes. A dissimilar-metal weld made between a low-alloy steel and an austenitic stainless steel, with an austenitic stainless steel filler metal, illustrates this problem. Chromium in steel has a greater affinity for carbon than does iron. When a carbon or low-alloy steel is welded with a steel filler metal containing a significant amount of chromium, carbon will diffuse from the base metal into the weld metal at temperatures above about 425 °C (800 °F). The diffusion rate is a function of temperature and exposure time, and it increases rather rapidly at 595 °C (1100 °F) and above. Carbon migration can take place during PWHT heat treatment or service at elevated temperatures.

Austenitic steel has a greater solubility for carbon than does ferritic steel. Therefore, carbon depletion in a carbon or lowalloy steel is greater when an austenitic stainless steel filler metal is used in preference to a ferritic steel filler metal. On the other hand, carbon migration is not as serious a problem when a nickel-chromium-iron filler metal is used. If carbon migration is extensive, it will be indicated by a lightly etching, low-carbon band in the carbon steel HAZ and a dark, high-carbon zone in the stainless steel weld metal, as seen in a transverse metallographic section. The extent of carbon migration during PWHT or elevated-temperature service should be determined during welding procedure qualification. It is known to influence long-time stress-rupture strength during elevated-temperature service. During cyclic temperature service, the HAZ will be subjected to varying shear stresses because of the differences in coefficients of thermal expansion of the base and weld metals. These stresses may produce fatigue failure in the decarburizing band next to the weld interface. Corrosion and Oxidation Resistance. The weld metal and both base metals have specific corrosion behaviors that must be considered by the designer in the initial selection of materials. For example, with dissimilar metal weldments, the formation of galvanic cells can cause corrosion of the most anodic metal or phase in the joint. Also, the weld metal is usually composed of several microstructural phases, and very localized cells between phases can result in galvanic corrosion at the microstructural level. To minimize galvanic corrosion, the composition of the weld metal can be adjusted to provide cathodic protection to the base metal that is most susceptible to galvanic attack. However, other design requirements should not be seriously compromised to do this. Instead, some other form of protection should be used.

A galvanic cell associated with a high-strength steel may promote hydrogen embrittlement in the HAZ of that steel if it is the cathode of the cell. Hydrogen embrittlement must be considered if the service temperature of the weldment will be in the range of -40 to 95 °C (-40 to 200 °F), and the weld will be in a highly stressed area of the assembly. Residual stresses developed in the weld zone are often sufficient to promote hydrogen embrittlement and stress-corrosion cracking. Chemical compositional differences in a dissimilar-metal weld can also cause high-temperature corrosion problems. Compositional variations at the interfaces between the different metals can result in selective oxidation when operating at high temperatures in air and formation of notches at these locations. Such notches are potential stress-raisers in the joint and can cause oxidation failure along the weld interface under cyclic thermal conditions. Specific Dissimilar Metal Combinations (Ref 33) The emphasis in the following section will be on joining carbon and low-alloy steels to various dissimilar materials (both ferrous and nonferrous) by arc welding. See other articles in this Section of the Handbook for additional information (for example, "Welding of Nickel Alloys" and "Welding of Copper Alloys"). Stainless Steels to Carbon or Low-Alloy Steels. Austenitic, ferritic, or martensitic stainless steels can be readily

fusion welded to carbon or low-alloy steels using a filler metal that can tolerate dilution by both base metals without formation of flaws in the joint. An austenitic stainless steel or nickel alloy filler metal is commonly used. The choice depends on the application and the service conditions. Generally, carbon or low-alloy steels are joined to 300-type stainless steels using austenitic filler metals, such as types E/ER 308, 309, 316, and 347. Such filler metals are used for service applications below 425 °C (800 °F). For higher temperatures or when high-temperature cyclic service is a concern, high-nickel filler metals such as ERNiCr-3 and ENiCrFe-3 are used. The coefficients of thermal expansion of high-nickel alloys approximate those of ferritic steels, and during cyclic-temperature service, the major differential expansion stresses are located primarily at the tough stainless steel/weld metal interface. Another advantage of a nickel-base weld metal is that it markedly reduces carbon migration from the ferritic steels to the weld metal. Extensive carbon migration into stainless steel weld metal weakens the HAZ of the carbon or low-alloy steel. The problem of cracking between ferritic steels, such as the 2.25Cr-1Mo steel described earlier, and austenitic stainless steels is a recurring problem in such operations as the petrochemical industry and electric utility generating plants. Cracking has been repeatedly observed in HAZs of the ferritic portions of dissimilar-metal welds, particularly those that have been subjected to a large number of thermal cycles and extended time at elevated temperatures around 565 °C (1050 °F). Such cracking must be minimized, since repair maintenance and non-productive downtime are costly. High-nickel filler metals have proven useful in such applications.

Buttering is also used to improve the weldability between ferritic and austenitic stainless steel joints. A thick overlay (13 mm, or

1 in., or greater) of a highly austenitic weld metal, such as type 309 stainless steel, can be deposited on carbon or 2

low-alloy steel components so that stainless steel/stainless steel welds can be made subsequently in the field. This technique has limited reliability for high-temperature cyclic service and has not been used extensively for material combinations other than ferritic steels, stainless steels, and some high-nickel alloys. A filler metal for joining a 4xx-series stainless steel to a carbon or low-alloy steel can be selected using the following general rules: •

•

•

FOR WELDING ONE HARDENABLE CHROMIUM STEEL TO ANOTHER WITH A HIGHER CHROMIUM CONTENT, FILLER METAL CONTAINING CHROMIUM EQUAL TO THAT OF EITHER STEEL MAY BE USED. FURTHERMORE, ANY FILLER METAL WHOSE CHROMIUM CONTENT LIES BETWEEN THESE LIMITS IS EQUALLY SATISFACTORY, PROVIDED THE WELDMENT IS PROPERLY HEAT-TREATED. A GENERAL RULE FOR WELDING ANY CHROMIUM STEEL TO ANY LOW-ALLOY STEEL IS TO USE A FILLER METAL THAT HAS THE SAME COMPOSITION AS THE LOW-ALLOY STEEL, PROVIDED THAT IT MEETS THE SERVICE REQUIREMENTS OF THE APPLICATION. WITH ANY LOW-ALLOY STEEL FILLER METAL, THE CHROMIUM THAT IS PICKED UP BY DILUTION WITH THE CHROMIUM STEEL BASE METAL MUST BE CONSIDERED. FOR WELDING ANY CHROMIUM STEEL TO A CARBON STEEL, CARBON STEEL FILLER METAL CAN BE USED. A CHROMIUM STEEL FILLER METAL CAN ALTERNATIVELY BE USED, BUT IT IS PREFERABLE TO USE A LESS HARDENABLE FILLER METAL.

When the properties of the HAZ in the chromium stainless steel are important, both the stainless steel and other steel can be buttered with type 309 or 310 austenitic stainless steel weld metal. An appropriate preheat or PWHT can be used to obtain desired properties in the buttered components. The stainless steel surfaces can then be welded together without preheat using a suitable austenitic stainless steel filler metal, such as type 308. Nickel-Base Alloys to Steels. Nickel alloys can easily be welded to steels using a suitable filler metal and proper

control of dilution. Nickel-base filler metals are generally used because of their good ductility and tolerance of dilution by iron. Suggested nickel-base filler metals for welding nickel alloys to steel or stainless steel are listed in Table 23.

TABLE 23 SUGGESTED FILLER METALS FOR WELDING NICKEL-BASE ALLOYS TO CARBON AND LOW-ALLOY STEELS NICKEL ALLOY UNS COMMON NO. DESIGNATION

FILLER METAL FORM

N02200

COMMERCIALLY PURE NICKEL

COVERED ELECTRODE BARE WIRE

N04400

MONEL 400

N05500 N05502 N06600

MONEL K-500 MONEL 502 INCONEL 600

COVERED ELECTRODE BARE WIRE

N08800

INCOLOY 800

N08825

INCOLOY 825

COVERED ELECTRODE BARE WIRE COVERED ELECTRODE BARE WIRE

FILLER METAL FOR WELDING TO STAINLESS CARBON OR STEEL LOW-ALLOY STEEL ENI-1, ENICRFE-2 ENI-1, ENICRFE-2, ENICRFE-3 ERNI-1, ERNICR-3 ERNI-1, ERNICR-3, ERNICRFE-6 ENICU-7, ENI-1 ENICRFE-2, ENICRFE-3 ERNI-1

ERNICR-3, ERNICRFE-6

ENICRFE-2, ENICRFE-3 ERNICR-3, ERNICRFE-6 ENICRMO-3

ENICRFE-2, ENICRFE-3 ERNICR-3, ERNICRFE-6

ERNICRMO-3

ERNICRMO-3

ENICRMO-3

N10665

HASTELLOY B-2

N10276

HASTELLOY C-276

N06455

HASTELLOY C-4

N06007

HASTELLOY G

COVERED ELECTRODE BARE WIRE COVERED ELECTRODE BARE WIRE COVERED ELECTRODE BARE WIRE COVERED ELECTRODE BARE WIRE

ENIMO-7

ENIMO-7

ERNIMO-7 ENICRMO-4

ERNIMO-7 ENICRMO-4

ERNICRMO-4 ENICRMO-4




ERNICRMO-1

ERNICRMO-1

Source: Ref 33

Sulfur and phosphorus in nickel and nickel alloys cause hot cracking. The melting techniques used to produce nickel and its alloys are designed to keep the content of these elements to low levels. By contrast, the sulfur and phosphorus contents in some steels are typically higher. Consequently, dilution should be carefully controlled when joining a steel to a nickel alloy with a nickel alloy filler metal, to avoid hot cracking in the weld metal. Most nickel-base weld metals can accept a substantial amount of iron dilution, but the dilution limit generally varies with the welding process. Weld metal deposited with nickel- or nickel-chromium-covered electrodes can tolerate up to approximately 40% iron dilution. With bare nickel or nickel-chromium filler metals, however, dilution should be kept to about 25%. Acceptable limits of iron dilution for nickel-copper weld metal vary, depending on the weld process. With SMAW, iron dilution of up to about 30% can be tolerated. Submerged arc weld metal should not be diluted by more than 25%. With the gas-shielding processes, nickel-copper weld metal is less tolerant of iron dilution, especially if the weld is to be thermally stress-relieved. The maximum limits for iron dilution in a welded joint are 10% when it will be used as-welded and 5% when it will be stress-relieved. A buttering layer of nickel or nickel-copper weld metal should be applied to the steel face in order to avoid exceeding these limits. Nickel-copper weld metal has a maximum dilution tolerance for chromium of about 8%. Consequently, nickel-copper filler metal should not be used to join nickel-copper alloys to stainless steels (see Table 23). Cobalt-Base Alloys to Steels. Metallurgically, cobalt-base alloys behave similarly to the high-temperature nickelchromium alloys with respect to welding. When joining a cobalt alloy to a stainless steel, a filler metal with a composition similar to that of the cobalt alloy is recommended. A nickel alloy filler metal may also be suitable for some applications. In any case, the filler metal selection, welding process, and welding procedure for the application should be established by suitable tests. Copper-Base Alloys to Steels. As described in the article "Welding of Copper Alloys" in this Volume, copper and

many copper-base alloys can be joined to carbon, low-alloy, and stainless steels by GTAW, GMAW, and SMAW. Iron dilution can be minimized by the use of appropriate welding procedures or by placement of a buttering layer of nickel on the steel. Tables 24 and 25 show combinations of dissimilar metals by GTAW and GMAW with the aid of copper-base or nickel-base filler metals.

TABLE 24 FILLER METALS AND PREHEAT AND INTERPASS TEMPERATURES USED IN GASTUNGSTEN ARC WELDING OF COPPERS AND COPPER-BASE ALLOYS TO STEELS METAL TO BE WELDED

LOW-CARBON STEEL

FILLER METALS (AND PREHEAT AND INTERPASS TEMPERATURES) FOR WELDING METAL IN COLUMN 1 TO(A): COPPERS PHOSPHOR ALUMINUM SILICON BRONZES BRONZES BRONZES ERCUAL-A2 OR ERCUSN-A (205 ERCUAL-A2 (150 ERCUAL-A2 (65 ERCU OR ERNI-3 °C, OR 400 °F) °C, OR 300 °F) °C, OR 150 °F

COPPER NICKELS ERCUAL-A2 OR ERNI-3 (65 °C, OR

MEDIUMCARBON STEEL HIGHCARBON STEEL LOW-ALLOY STEEL STAINLESS STEEL

(540 °C, OR 1000 °F) ERCUAL-A2 OR ERCU OR ERNI-3 (540 °C, OR 1000 °F) ERCUAL-A2 OR ERCU OR ERNI-3 (540 °C, OR 1000 °F) ERCUAL-A2 OR ERCU OR ERNI-3 (540 °C, OR 1000 °F) ERCUAL-A2 OR ERCU OR ERNI-3 (540 °C, OR 1000 °F)

ERCUSN-A (205 °C, OR 400 °F)

ERCUAL-A2 (205 °C, OR 400 °F)

ERCUSN-A (260 °C, OR 500 °F)

ERCUAL-A2 (260 °C, OR 500 °F)

ERCUSN-A (260 °C, OR 500 °F)

ERCUAL-A2 (260 °C, OR 500 °F)

ERCUSN-A (205 °C, OR 400 °F)

ERCUAL-A2 (65 °C, OR 150 °F MAX)

MAX) ERCUAL-A2 (65 °C, OR 150 °F MAX) ERCUAL-A2 (205 °C, OR 400 °F) ERCUAL-A2 (205 °C, OR 400 °F) ERCUAL-A2 (65 °C, OR 150 °F MAX)

150 °F MAX) ERCUAL-A2 OR ERNI-3 (65 °C, OR 150 °F MAX) ERCUAL-A2 OR ERNI-3 (65 °C, OR 150 °F MAX) ERCUAL-A2 OR ERNI-3 (65 °C, OR 150 °F MAX) ERCUAL-A2 OR ERNI-3 (65 °C, OR 150 °F MAX)

(A) FILLER-METAL SELECTIONS SHOWN IN TABLE ARE BASED ON WELDABILITY, EXCEPT WHERE MECHANICAL PROPERTIES ARE USUALLY MORE IMPORTANT. PREHEATING IS ORDINARILY USED ONLY WHEN AT LEAST ONE MEMBER IS THICKER THAN ABOUT 3.2 MM (

1 8

IN.) OR IS HIGHLY CONDUCTIVE (SEE TEXT). PREHEAT AND INTERPASS

TEMPERATURES ARE SUBJECT TO ADJUSTMENT ON THE BASIS OF THE SIZE AND SHAPE OF THE WELDMENT.

TABLE 25 FILLER METALS AND PREHEAT AND INTERPASS TEMPERATURES USED IN GAS-METAL ARC WELDING OF COPPERS AND COPPER-BASE ALLOYS TO STEELS METAL TO BE WELDED

LOW-CARBON STEEL MEDIUMCARBON STEEL HIGH-CARBON STEEL LOW-ALLOY STEEL STAINLESS STEEL

ELECTRODES (AND PREHEAT AND INTERPASS TEMPERATURES) FOR WELDING METAL IN COLUMN 1 TO(A): COPPERS LOW-ZINC HIGH-ZINC PHOSPHOR BRASSES BRASSES, BRONZES TIN BRASSES, SPECIAL BRASSES ERCUAL-A2 OR ERCUSN-A (315 °C, ERCUAL-A2 ERCUSN-A (205 ERCU OR ERNI-3 (540 OR 600 °F) (260 °C, OR 500 °C, OR 400 °F) °C, OR 1000 °F) °F) ERCUAL-A2 OR ERCUAL-A2 (315 ERCUAL-A2 ERCUSN-A (205 ERCU OR ERNI-3 (540 °C, OR 600 °F) (260 °C, OR 500 °C, OR 400 °F) °C, OR 1000 °F) °F) ERCUAL-A2 OR ERCUAL-A2 (315 ERCUAL-A2 ERCUSN-A (260 ERCU OR ERNI-3 (540 °C, OR 600 °F) (260 °C, OR 500 °C, OR 500 °F) °C, OR 1000 °F) °F) ERCUAL-A2 OR ERCUAL-A2 (315 ERCUAL-A2 ERCUSN-A (260 ERCU OR ERNI-3 (540 °C, OR 600 °F) (315 °C, OR 600 °C, OR 500 °F) °C, OR 1000 °F) °F) ERCUAL-A2 OR ERCUAL-A2 OR ERCUAL-A2 ERCUSN-A (205 ERCU OR ERNI-3 ERCUSN-A (315 °C, (315 °C, OR 600 °C, OR 400 °F) (540 °C, OR 1000 °F) OR 600 °F) °F)

ALUMINUM BRONZES

SILICON BRONZES

COPPER NICKELS

ERCUAL-A2 (150 °C, OR 300 °F)

ERCUAL-A2 (260 °C, OR 500 °F)

ERCUAL-A2 (65 °C, OR 150 °F MAX) ERCUAL-A2 (65 °C, OR 150 °F MAX) ERCUAL-A2 (205 °C, OR 400 °F)

ERCUAL-A2 (260 °C, OR 500 °F)

ERCUAL-A2 (205 °C, OR 400 °F)



ERCUAL-A2OR ERNI-3 (65 °C, OR 150 °F MAX) ERCUAL-A2 OR ERNI-3 (65 °C, OR 150 °F MAX) ERCUAL-A2 OR ERNI-3 (65 °C, OR 150 °F MAX) ERCUAL-A2, OR ERNI-3 (65 °C, OR 150 °F MAX) ERCUAL-A2 OR ERNI-3 (65 °C, OR 150 °F MAX)

ERCUAL-A2 (205 °C, OR 400 °F)

(A) ELECTRODE SELECTIONS IN TABLE ARE BASED ON WELDABILITY, EXCEPT WHERE MECHANICAL PROPERTIES ARE USUALLY MORE IMPORTANT. PREHEATING USUALLY IS USED ONLY WHEN AT LEAST ONE MEMBER IS THICKER THAN 3.2 MM (

1 8

IN.) OR

IS HIGHLY CONDUCTIVE (SEE TEXT). PREHEAT AND INTERPASS TEMPERATURES ARE SUBJECT TO ADJUSTMENT BASED ON SIZE AND SHAPE OF WELDMENT.

Aluminum-Base Alloys to Steels. With respect to fusion welding, iron and aluminum are not compatible metals. Their

melting temperatures differ greatly: 660 °C (1220 °F) for aluminum versus 1538 °C (2800 °F) for iron. Both metals have almost no solubility for the other in the solid state, especially iron in aluminum, and several brittle intermetallic phases can form (FeAl2, Fe2Al5, or FeAl3). Consequently, fusion welds joining iron and aluminum are brittle. In addition, high welding stresses would be expected because of the significant differences in the thermal expansion coefficients, thermal conductivities, and specific heats of iron and aluminum. Aluminum can be joined to carbon or stainless steel by brazing, solid-state welding or by high-energy processes such as electron-beam welding.


32. D.J. KOTECKI AND T.A. SIEWERT, WRC-1992 CONSTITUTION DIAGRAM FOR STAINLESS STEEL WELD METALS: A MODIFICATION OF THE WRC-1988 DIAGRAM, WELD. J., MAY 1992, P 171S-178S 33. DISSIMILAR METALS, WELDING HANDBOOK, 7TH ED., VOL 4, AMERICAN WELDING SOCIETY, 1982, P 514-547 34. DISSIMILAR-WELD FAILURE ANALYSIS AND DEVELOPMENT PROGRAM, VOL 1-4, REPORT CS-4252, ELECTRIC POWER RESEARCH INSTITUTE, PALO ALTO, CA Hardfacing, Weld Cladding, and Dissimilar Metal Joining J.R. Davis, Davis & Associates

References

1. K.C. ANTONY ET AL., HARDFACING, WELDING, BRAZING, AND SOLDERING, 9TH ED., VOL 6, METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 771-793 2. P. CROOK AND H.N. FARMER, FRICTION AND WEAR OF HARDFACING ALLOYS, FRICTION, LUBRICATION, AND WEAR TECHNOLOGY, VOL 18, ASM HANDBOOK, ASM INTERNATIONAL, 1992, P 758-765 3. D.K. SUBRAMANYAM, A.E. SWANSIGER, AND H.S. AVERY, AUSTENITIC MANGANESE STEELS, PROPERTIES AND SELECTION: IRONS, STEELS, AND HIGH-PERFORMANCE ALLOYS, VOL 1, ASM HANDBOOK, ASM INTERNATIONAL, 1990, P 822-840 4. H.S. AVERY AND H.J. CHAPIN, WELD J., VOL 31, 1952, P 917 5. R.D. ZORDAN, REPORT 10890, CABOT CORPORATION, KOKOMO, IN, 1982 6. H. OCKEN AND A.K. VELAN, ED., "LABORATORY EVALUATIONS OF COBALT-FREE, NICKELBASED HARDFACING ALLOYS FOR NUCLEAR APPLICATIONS," REPORT NP-4993, PROJECT 1935, ELECTRIC POWER RESEARCH INSTITUTE, PALO ALTO, CA, MARCH 1987 7. P.J. HOFMANN, B.C. FRIEDRICH, AND H. OCKEN, ED., "LABORATORY EVALUATIONS OF IRON-BASED HARDFACING ALLOYS," REPORT NP-5874, PROJECT 1935-11, ELECTRIC POWER RESEARCH INSTITUTE, PALO ALTO, CA, JUNE 1988 8. WELD DESIGN AND FABRICATION, DATA SHEET 420, WELD. DES. FABR., VOL 48 (NO. 8), 1975, P 50 9. K.M. KULKARNI AND V. ANAND, METAL POWDERS USED FOR HARDFACING, POWDER METALLURGY, 9TH ED., VOL 7, ASM HANDBOOK, AMERICAN SOCIETY FOR METALS, 1984, P 823-836 10. K.P. COOPER, LASER SURFACE PROCESSING, FRICTION, LUBRICATION, AND WEAR TECHNOLOGY, VOL 18, ASM HANDBOOK, ASM INTERNATIONAL, 1992, P 861-872 11. A.G. BLAKE, ET AL., "LASER COATING TECHNOLOGY: A COMMERCIAL REALITY," BULLETIN OF QUANTUM LASER CORPORATION, EDISON, NJ

12. J.D. AYERS AND D.S. GNANAMUTHU, HARDENING OF METAL SURFACES BY LASER PROCESSING, WELDING, BRAZING, AND SOLDERING, 9TH ED., VOL 6, ASM HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 793-803 13. W.M. STEEN, PROC. NATO ADVANCED STUDY INSTITUTE ON LASER SURFACE TREATMENT OF METALS, C.W. DRAPER AND P. MAZZOLDI, ED., MARTINUS NIJHOFF, 1986, P 369 14. J. MAZUMDER AND J. SINGH, PROC. NATO ADVANCED STUDY INSTITUTE ON LASER SURFACE TREATMENT OF METALS, C.W. DRAPER AND P. MAZZOLDI, ED., MARTINUS NIJHOFF, 1986, P 297 15. A. BELMONDO AND M. CASTAGNA, SOURCEBOOK ON APPLICATIONS OF THE LASER IN METALWORKING, E. METZBOWER, ED., AMERICAN SOCIETY FOR METALS, 1981, P 310 16. G.J. BRUCK, J. MET., VOL 39 (NO. 2), 1987, P 10 17. A.V. LA ROCCA, PROC. NATO ADVANCED STUDY INSTITUTE ON LASER SURFACE TREATMENT OF METALS, C.W. DRAPER AND P. MAZZOLDI, ED., MARTINUS NIJHOFF, 1986, P 521 18. R.M. MACINTYRE, PROC. NATO ADVANCED STUDY INSTITUTE ON LASER SURFACE TREATMENT OF METALS, C.W. DRAPER AND P. MAZZOLDI, ED., MARTINUS NIJHOFF, 1986, P 545 19. B.A. KUSHNER AND E.R. NOVINSKI, THERMAL SPRAY COATINGS, FRICTION, LUBRICATION, AND WEAR TECHNOLOGY, VOL 18, ASM HANDBOOK, ASM INTERNATIONAL, 1992, P 829-833 20. R.H. UNGER, THERMAL SPRAY COATINGS, CORROSION, 9TH ED., VOL 13, ASM HANDBOOK, ASM INTERNATIONAL, 1987, P 459-462 21. M.L. THORPE, THERMAL SPRAY: INDUSTRY IN TRANSITION, ADV. MATER. PROCESS., VOL 143 (NO. 5), 1993, P 50-56 22. G. IRONS, "HIGHER VELOCITY THERMAL SPRAY PROCESSES PRODUCE BETTER AIRCRAFT ENGINE COATINGS," SAE TECHNICAL PAPER 920947, PRESENTED AT THE 28TH ANNUAL AEROSPACE/AIRLINE PLATING AND METAL FINISHING FORUM AND EXPOSITION (SAN DIEGO, CA), APRIL 1992 23. F.N. LONGO, PROCESSING AND DESIGN, LESSON 3, COURSE 51, THERMAL SPRAY TECHNOLOGY, MATERIALS ENGINEERING INSTITUTE, ASM INTERNATIONAL, 1992 24. M.L. THORPE, THERMAL SPRAYING BECOMES A DESIGN TOOL, MACH. DES., 24 NOV 1983 25. P.O. BUZBY AND J. NIKITICH, HVOF THERMAL SPRAYING OF NITRIDED PARTS, ADV. MATER. PROCESS., DEC 1991, P 35-36 26. R.D. THOMAS ET AL., WELD OVERLAYS, WELDING, BRAZING, AND SOLDERING, 9TH ED., VOL 6, METALS HANDBOOK, AMERICAN SOCIETY FOR METALS, 1983, P 804-819 27. D.L. OLSON, PREDICTION OF AUSTENITIC WELD METAL MICROSTRUCTURE AND PROPERTIES, WELD. J., VOL 64 (NO. 10), 1985, P 2815-2955 28. A.L. SCHAEFFLER, CONSTITUTION DIAGRAM FOR STAINLESS STEEL WELD METAL, MET. PROG., VOL 56 (NO. 11), 1949, P 680-680B 29. W.T. DELONG, FERRITE IN AUSTENITIC STAINLESS STEEL WELD METAL, WELD. J., VOL 53 (NO. 7), 1974, P 2735-2865 30. T.A. SIEWERT, C.N. MCCOWAN, AND D.L. OLSON, FERRITE NUMBER PREDICTION TO 100 FN IN STAINLESS STEEL WELD METAL, WELD. J., VOL 67 (NO. 12), 1988, P 2895-2985 31. D.J. KOTECKI, "VERIFICATION OF THE NBS-CSM FERRITE DIAGRAM," INTERNATIONAL INSTITUTE OF WELDING DOCUMENT II-C-834-88, 1988 32. D.J. KOTECKI AND T.A. SIEWERT, WRC-1992 CONSTITUTION DIAGRAM FOR STAINLESS STEEL WELD METALS: A MODIFICATION OF THE WRC-1988 DIAGRAM, WELD. J., MAY 1992, P 171S-178S 33. DISSIMILAR METALS, WELDING HANDBOOK, 7TH ED., VOL 4, AMERICAN WELDING SOCIETY, 1982, P 514-547 34. DISSIMILAR-WELD FAILURE ANALYSIS AND DEVELOPMENT PROGRAM, VOL 1-4, REPORT

CS-4252, ELECTRIC POWER RESEARCH INSTITUTE, PALO ALTO, CA

Procedure Development and Practice Considerations for Resistance Welding

Introduction RESISTANCE WELDING (RW) encompasses a group of processes in which the heat for welding is generated by the resistance to the flow of electrical current through the parts being joined. These processes are most commonly used to weld two overlapping sheets or plates that may have different thicknesses. A pair of electrodes clamp the workpieces under pressure to provide good electrical contact and to contain the molten metal in the joint. The electrodes conduct electrical current to the joint; resistance to the flow of current heats the faying surfaces, forming a weld (Ref 1). The required heat generated through the electrical resistance of the two bodies to be joined is given by

H = I2RT

(EQ 1)

where H is the heat generated in joules; I is the current in amperes; R is the resistance in ohms; and t is the time of current flow, in seconds. The actual temperature rise at the joint depends on the interrelationship of the three component variables. These variables are affected by: • • •

COOLING EFFECT OF THE ELECTRODE EFFECT OF FOLLOW UP ON RESISTANCE EFFECT OF SHUNTING

Pressure is applied even before the current is activated to ensure a good bond. Accurate control and timing of the electric current and application of pressure is an essential feature of resistance welding. The strength of the bond also depends on the cleanliness of the mating surfaces, although the process is more tolerant of the presence of oxide layers and contaminants than other welding processes. Resistance welding processes require specialized machinery, but they have the major advantages of not requiring filler metals, shielding gases, or flux. The three major resistance welding processes are resistance spot welding (RSW), resistance seam welding (RSEW), and projection welding (PW). Flash welding (FW) is also classified as a resistance welding process, but it is unique in that the heat for welding is created at the faying surfaces of the joint not only by resistance to the flow of electric current, but also by arcs across the interface. Upset welding (UW) differs from flash welding in that no arc is generated by the process. Capacitor discharge stud welding incorporates a stud which serves as a weld member and is a process similar to flash welding. High-frequency resistance welding (HFRW) concentrates current of the surface of the part to heat the base metal and cause the coalescence of metal. This article addresses procedure development and considerations for using these processes to join specific types of materials. Additional information about these processes can be found in the Section "Fusion Welding Processes" in this Volume. Acknowledgements Special thanks are due to David M. Beneteau of CenterLine (Windsor) Limited, who offered suggestions about revisions for portions of this article. His participation was arranged through the Resistance Welder Manufacturers Association.

Reference

1. L.P. CONNER, ED., WELDING HANDBOOK, 8TH ED., VOL 1, AWS, 1987, P 16 Procedure Development and Practice Considerations for Resistance Welding

Resistance Spot Welding

Resistance spot welding is the most common form of resistance welding. It is a process in which laying surfaces are joined in one or more overlapping spots by the heat generated by resistance to the flow of electric current through workpieces that are held together under force by electrodes. The contacting surfaces in the region of current concentration are heated by a short-time pulse of low-voltage, high-amperage current to form a fused nugget of weld metal. When the flow of current ceases, the electrode force is maintained while the weld metal rapidly cools and solidifies. The electrodes are retracted after each weld, which usually is completed in a fraction of a second. The size and shape of the individually formed welds are limited primarily by the size and contour of the electrode faces. The weld nugget forms at the faying surfaces (Fig. 1), but does not extend completely to the outer surfaces. In a cross section, the nugget in a properly formed spot weld is obround or oval in shape; in plan view, it has the same shape as the electrode face (which usually is round) and approximately the same size. The spots should be at a sufficient distance from the edge of the workpiece (edge distance) so that there is enough base metal to withstand the electrode force and to ensure that the local distortion during welding does not allow expulsion of metal from the weld. Additional information is available in the article "Resistance Spot Welding" in this Volume.

FIG. 1 CUTAWAY VIEW SHOWING LOCATION AND SHAPE OF WELD NUGGET RELATIVE TO THE INNER AND OUTER SURFACES OF WORKPIECES IN A RESISTANCE SPOT WELDING SETUP

RSW Applications Weldability and the thickness of the metals joined are key factors in determining whether two materials can be resistance spot welded. Material Selection (Ref 2). Practically all combinations of ductile metals and alloys can be spot welded (Table 1).

Some, like copper to aluminum, and aluminum to magnesium, form alloys having little strength. Others, such as zinc and some of the high-chromium alloys, experience grain growth even during a very short welding period. Although welding times are relatively short, the material is adversely affected by the heat conducted through the sheet to the point that the physical properties such as the strength and corrosion resistance of the material surface are typically reduced in value. The short welding period is also necessary in stainless steel to prevent carbide precipitation when the carbon content is high enough to permit it. High-carbon steels weld readily, but the weld will be in the full-hardened state and will require subsequent heat treatment. This can be accomplished automatically by a control that applies preheating, postheating, or both as part of the welding cycle. To ensure that the control can apply the compensating quench or temper cycle, the electrodes must open immediately after the cycle is complete. Excessive hold time will quench the joint, negating the benefits associated with this control function.

TABLE 1 RELATIVE WELDABILITY RATINGS OF SELECTED METAL AND ALLOY COMBINATIONS USED FOR RESISTANCE SPOT WELDING OPERATIONS A, excellent; B, good; C, fair; D, poor; E, very poor; F, impractical METALS

ALUMIN UM

STAINL ESS STEEL

BRA SS

COPP ER

GALVANI ZED IRON

STE EL

LEA D

MON EL

NICK EL

NICHRO ME

TINPLA TE

ZIN C

PHOSPH OR BRONZE

NICK EL SILV ER

TERNEPL ATE

ALUMINU M STAINLES S STEEL BRASS COPPER GALVANI ZED IRON STEEL LEAD MONEL NICKEL NICHROM E TINPLATE ZINC PHOSPHO R BRONZE NICKEL SILVER TERNEPL ATE

B

E

D

E

C

D

E

D

D

D

C

C

C

F

C

F

A

E

E

B

A

F

C

C

C

B

F

D

D

B

D E C

E E B

C D D

D F E

D E B

D E B

F E D

C D C

C D C

C D C

D E B

E E C

C C D

C C E

D E B

D E D D D

A F C C C

D F C C C

E E D D D

B D C C C

A E C C C

E C E E E

C E A B B

C E B A B

C E B B A

B ... C C C

F C F F F

C E C C D

D E B B B

A D C C C

C C C

B E D

D E C

E E C

B C D

B F C

... C E

C F C

C F C

C F D

C C D

C C D

D D B

D F B

C C D

F

D

C

C

E

D

E

B

B

B

D

F

B

A

D

C

B

D

E

B

A

D

C

C

C

C

C

D

D

B

Source: Ref 3

Zinc-alloy die castings can be welded with little loss of strength, but the ductility is reduced. Free-turning Bessemer screw-machine steel frequently refuses to weld, or results in brittle welds, and hence its use for welded parts should be avoided. Usually, however, when clean, properly prepared metal parts are spot welded, the strength of the welds is perfectly satisfactory. Table 1 lists combinations that have been successfully welded. Many other combinations may also be satisfactorily welded. Copper and silver are difficult to weld, but they may be welded by the use of low-conductivity electrodes (for example, refractory alloys such as Elkonite and Trodaloy). Copper alloys are more commonly resistance brazed. In many applications, the weldability of copper is increased if it has a tinned surface. Thickness Parameters (Ref 4). Parts of widely different thicknesses may be spot welded. For example, 0.4 mm (

1 64

in.) material can be welded to a 150 mm (6 in.) thick piece and would require only slightly more power input and pressure than to weld two pieces of 0.4 mm (

1 in.) material, as it is not necessary to heat the 150 mm (6 in.) piece through, and 64

the push-up is obtained on the thinner piece. The thickness of the thinner component is referred to as the governing metal thickness because it dictates the heat input. However, the thermal mass of the larger component may make RSW applications impractical because the thick section may not reach fusion temperature before the thin piece is melted. The limit to welding pieces of equal thickness with an uninterrupted flow of current seems to be approximately 3.2 mm (

1 8

in.). Greater thicknesses may be welded by a technique known as pulsation welding. In the pulsation welding method, the current is turned off during the weld at intervals that allow the electrodes to cool the electrode/workpiece surface. This ensures that the heat will be generated as desired, primarily at the faying surface. Sheet Steel Limitations. Spot welded lap joints are widely used in joining sheet steel up to about 3.2 mm (

(Table 2) and are used occasionally in joining steel 6.4 mm (

1 in.) thick 8

1 in.) or more in thickness, t. Thicknesses of 25 mm (1 in.) 4

or more have been joined by spot welding, but this requires special equipment and would not ordinarily be economical. Many assemblies of two or more sheet metal stampings that do not require gastight or liquid-tight joints can be more economically joined by high-speed RSW than by mechanical methods.

TABLE 2 MINIMUM DIMENSIONS FOR SINGLE-IMPULSE SPOT WELDS IN LOW-CARBON STEEL

STOCK THICKNESS MINIMUM WELD SPACING, S mm in. mm in. 0.25 0.010 6.3 0.250 0.53 0.021 9.5 0.375 0.80 0.031 13 0.500 1.0 0.040 19 0.750 1.3 0.050 22 0.875 1.6 0.063 27 1.062

MINIMUM OVERLAP OF MATERIAL, L in. mm 9.5 0.375 11 0.433 11 0.433 13 0.500 14 0.551 16 0.625

2.0 2.4 2.8 3.2

0.078 0.094 0.109 0.125

35 41 46 51

1.375 1.625 1.813 2.000

17 19 21 22

0.669 0.750 0.827 0.875

Source: Ref 5

For aluminum and magnesium welds, minimum spot spacing, S, should equal 8t. A spot spacing value of 16t would be even better if the strength requirements of the joint permit such spacing (Ref 5). Joint Design Recommendations. Like other welds, resistance spot welds are subject to shrinkage on cooling. To

minimize distortion, welds should be positioned to balance shrinkage forces (Fig. 2).

FIG. 2 PLACEMENT OF WELDS TO JOIN TWO SHEETS OF METAL IN A T-JOINT. (A) CONFIGURATION NOT RECOMMENDED DUE TO UNBALANCED SHRINKAGE FORCES. (B) PREFERRED SETUP THAT RESULTS IN BALANCED FORCES AND MINIMUM DISTORTION

Distortion can also result from electrode skidding. This is related to stress introduced at the time of electrode impact and machine deflection under high electrode force. The components must be designed to provide methods of accommodating variations (for example, by incorporating slip joints and minimizing material stackups). Even when welding of multiple spot thicknesses is feasible, it can be more troublesome than simple welds of two sheet thicknesses. Scalloped edges permit alternate welding of spots in two thicknesses only (Fig. 3). However, scalloped edges are rarely used and the current trend is toward triple-metal-thickness welding as a method of reducing capital cost.

FIG. 3 PLACEMENT OF WELDS TO JOIN THREE SHEETS OF METAL IN A T-JOINT. (A) CONFIGURATION THAT CAN BE USED TO JOIN THREE STOCK THICKNESSES. (B) RECOMMENDED CONFIGURATION INCORPORATING

SCALLOPED EDGES TO PERMIT ALTERNATE WELDING OF SPOTS IN TWO THICKNESSES ONLY

Spot-welded assemblies should be designed so that joint areas are readily accessible to the electrode and simple, inexpensive electrodes can be used. When special electrode shapes are required, every effort should be made to avoid extremes of shape and to provide as accessible a joint as possible. Deep, narrow joint areas should be avoided. Figure 4 provides guidelines for accessibility of joints.

FIG. 4 GRADUAL IMPROVEMENTS IN JOINT DESIGN THAT PROVIDE OPTIMUM ACCESS TO SELECTED JOINT CONFIGURATIONS USING STANDARD ELECTRODES

Weld Strength in RSW Joints. The strength of a single spot weld in shear depends on the cross-sectional area of the nugget in the plane of the faying surfaces.

Lap joints may fail either by shear through the nugget, or by tearing of the base metal adjacent to the weld nugget. Low weld strengths are typically associated with weld nugget shear failure, and high strengths are typically associated with base metal tearing. A minimum nugget diameter that is unique to the type of base material, surface condition, and metal gage used, is required in order to obtain failure by base metal tearing. Increasing the nugget diameter above this minimum value can produce an increase in the weld strength. Figure 5 shows the slight increase in strength values obtained for low-carbon steel as a result of increases in nugget size.

FIG. 5 PLOT OF TENSILE SHEAR STRENGTH VERSUS WELD NUGGET DIAMETER AS A FUNCTION OF SHEET THICKNESS, T, FOR 1008 LOW-CARBON STEEL

The strength of multiple spot-welded joints depends on material thickness, spot weld spacing, and weld pattern. The spacing between adjacent spot welds may alter the joint weld strength due to current shunting through previous welds. As the spacing between adjacent spot welds decreases, the joint shear strength may decrease. Table 3 summarizes the relationship between resistance spot welding variables and joint strength for uncoated low-carbon steel.

TABLE 3 RECOMMENDED PARAMETERS FOR RESISTANCE SPOT WELDING OF UNCOATED LOW-CARBON STEEL SHEET PRODUCED FOR AUTOMOTIVE APPLICATIONS Data are average values. SHEET THICKNESS

MM 0.51 0.64 0.76 0.89 1.02 1.14 1.27 1.40 1.52 1.78 2.03 2.29 2.67 3.05

IN. 0.020 0.025 0.030 0.035 0.040 0.045 0.050 0.055 0.060 0.070 0.080 0.090 0.105 0.120

FORCE(C)

ELECTRODE FACE DIAMETER MM IN. 4.78 0.188 4.78 0.188 6.35 0.250 6.35 0.250 6.35 0.250 6.35 0.250 7.92 0.312 7.92 0.312 7.92 0.312 7.92 0.312 7.92 0.312 9.53 0.375 9.53 0.375 9.53 0.375

SHAPE(A)

BEVEL ANGLE(B)

E,A,B E,A,B E,A,B E,A,B E,A,B E,A,B E,A,B E,A,B E,A,B E,A,B E,A,B E,A,B E,A,B E,A,B

45 ° 45 ° 45 ° 45 ° 45 ° 45 ° 30° 30° 30° 30° 30° 30° 30° 30°

KN 1.8 2.0 2.2 2.7 3.1 3.4 3.6 4.0 4.5 5.4 6.3 7.2 8.1 9.4

LBF 400 450 500 600 700 750 800 900 1000 1200 1400 1600 1800 2100

WELD TIME, CYCLES (60 HZ)

WELDING CURRENT, A


7 8 9 9 10 11 12 13 14 16 18 20 23 26

8,500 9,500 10,500 11,500 12,500 13,000 13,500 14,000 15,000 16,000 17,000 18,000 19,500 21,000

MM 11.2 11.9 11.9 13.5 13.5 15.0 15.0 16.0 16.0 16.8 18.3 19.8 21.3 22.4

IN. 0.44 0.47 0.47 0.53 0.53 0.59 0.59 0.63 0.63 0.66 0.72 0.78 0.84 0.88

MINIMUM SPACING(D) TWO STACK MM IN. 9.7 0.38 15.8 0.62 15.8 0.62 19.1 0.75 19.1 0.75 24.9 0.94 24.9 0.94 26.9 1.06 26.9 1.06 30.0 1.18 35.1 1.38 39.6 1.56 42.7 1.68 46.0 1.81

WELD THREE STACK MM IN. 15.8 0.62 22.4 0.88 22.4 0.88 26.9 1.06 26.9 1.06 30.0 1.18 30.0 1.18 33.3 1.31 33.3 1.31 38.1 1.50 40.6 1.60 47.8 1.88 50.8 2.00 63.5 2.50

MINIMUM SHEAR STRENGTH

BUTTON DIAMETER

KN 1.43 2.02 2.58 3.36 4.14 5.15 6.05 7.53 8.29 10.3 12.1 15.5 18.6 22.4

MM 2.5 3.1 3.6 4.1 4.6 4.8 5.1 5.3 5.8 6.4 6.6 6.9 7.1 7.6

LBF 320 450 575 750 925 1150 1350 1680 1850 2300 2700 3450 4150 5000

Source: Ref 6

(A) (B) (C) (D)

SHAPE DEFINITIONS: E, TRUNCATED CONE; A, POINTED; B, 75 MM (3 IN.) RADIUS. APPLIES TO TRUNCATED CONE ELECTRODES ONLY AND IS MEASURED FROM THE PLANE OF THE ELECTRODE FACE. FOR INTERMEDIATE THICKNESSES, FORCE AND WELD TIME MAY BE INTERPOLATED. MINIMUM WELD SPACING IS MEASURED FROM CENTERLINE TO CENTERLINE.

IN. 0.10 0.12 0.14 0.16 0.18 0.19 0.20 0.21 0.23 0.25 0.26 0.27 0.28 0.30

Power Requirements. Nearly all RSW of low-carbon steel is direct-energy welding, in which single- or three-phase 60-

cycle alternating current, drawn ordinarily from 220 or 440 V implant power lines and stepped down to about 2 to 20 V, is fed directly to the electrodes as each weld is made. Direct and Series Welding Spot welds are classified as either direct or series welds. Direct Single-Spot Welding. Single-spot welds are usually made by direct welding. Figure 6 shows schematically three

arrangements used for making this type of weld; these arrangements may be modified to meet special requirements. In all of the arrangements shown, one transformer secondary circuit makes one spot weld.

FIG. 6 RELATION OF WORK METAL TO ELECTRODES WHEN MAKING SINGLE-SPOT WELDS BY DIRECT WELDING

The simplest and most common arrangement in which two workpieces are sandwiched between opposing upper and lower electrodes, is shown in Fig. 6(a). In Fig. 6(b), a conductive plate or mandrel having a large contacting surface is used as the lower electrode; this reduces marking on the lower workpiece and conducts heat away from the weld more rapidly and may be necessary because of the shape of the workpiece. In the arrangement in Fig. 6(c), a conductive plate or mandrel beneath the lower workpiece is used for the same purposes but in conjunction with a second upper electrode. Direct Multiple-Spot Welding. Two arrangements of the secondary circuit for making two or more spot welds simultaneously by direct welding are shown in Fig. 7(a) and 7(b). In direct multiple-spot welding (as well as in series multiple-spot welding, described below), tip contour and surface condition must be the same for each electrode. Also, the force exerted by all the electrodes on the workpieces must be equal, regardless of inequalities in work-metal thickness. The force can be equalized by using a spring-loaded electrode holder or a hydraulic equalizing system. The use of a conductive plate or mandrel, as in Fig. 7(b), minimizes weld marks on the lower workpiece.

FIG. 7 RELATION OF WORK METAL TO ELECTRODES WHEN MAKING MULTIPLE-SPOT WELDS. (A) AND (B) DIRECT WELDING. (C), (D), AND (E) SERIES WELDING

Series Multiple-Spot Welding. Three arrangements for making a number of spot welds simultaneously by series

welding are shown in Fig. 7(c), 7(d), and 7(e). In Fig. 7(d), each of the two transformer secondary circuits makes two spot welds. A portion of the current bypasses the weld nuggets through the upper workpiece.

Push-Pull Welding. Figure 7(e) is commonly referred to as push-pull welding. The advantage to this process is that the

secondary loop area is quite small. This is common for components such as floor pans where a normal welding unit would have a throat several feet deep.


2. J.F. YOUNG AND R.S. SHANE, ED., MATERIALS AND PROCESSES, PART B: PROCESSES, 3RD ED., MARCEL DEKKER, 1985, P 11257-1126 3. J.G. BRALLA, ED., HANDBOOK OF PRODUCT DESIGN FOR MANUFACTURING, MCGRAW-HILL, 1986, P 7-35 4. J.F. YOUNG AND R.S. SHANE, ED., MATERIALS AND PROCESSES, PART B: PROCESSES, 3RD ED., MARCEL DEKKER, 1985, P 1127

5. J.G. BRALLA, ED., HANDBOOK OF PRODUCT DESIGN FOR MANUFACTURING, MCGRAW-HILL, 1986, P 7-38 6. R.L. O'BRIEN, ED., WELDING HANDBOOK, 8TH ED., VOL 2, AWS, 1991, P 552 Procedure Development and Practice Considerations for Resistance Welding

Resistance Seam Welding Resistance seam welding is a process in which heat caused by resistance to the flow of electric current in the work metal is combined with pressure to produce a welded seam. This seam, consisting of a series of overlapping spot welds, is normally gastight or liquid-tight. Two rotating, circular electrodes (electrode wheels), or one circular and one bar-type electrode, are used for transmitting the current to the work metal. When two electrode wheels are used, one or both wheels are driven either by means of a gear-driven shaft or by a knurl or friction drive that contacts the peripheral surface of the electrode wheel. The series of spot welds is made without retracting the electrode wheels or releasing the electrode force between spots, although the electrode wheels may advance either continuously or intermittently. Additional information is available in the article "Resistance Seam Welding" in this Volume. Seam Welding Machines A seam welding machine is similar in construction to a spot welding machine, except that one or two electrode wheels are substituted for the spot welding electrodes. Generally, seam. Welding is done in a press-type resistance welding machine. There are four basic types of RSEW machines: • •

•

•

CIRCULAR: AXIS OF ROTATION OF THE ELECTRODE WHEELS IS AT RIGHT ANGLES TO THE FRONT OF THE MACHINE (FIG. 8A). LONGITUDINAL: AXIS OF ROTATION OF THE ELECTRODE WHEELS IS PARALLEL TO THE FRONT OF THE MACHINE (FIG. 8B), AND THROAT DEPTH IS TYPICALLY 305 TO 915 MM (12 TO 36 IN.). UNIVERSAL: A SWIVEL-TYPE HEAD AND INTER-CHANGEABLE LOWER ARMS ALLOW THE AXIS OF ROTATION OF THE ELECTRODE WHEELS TO BE SET EITHER AT RIGHT ANGLES OR PARALLEL TO THE FRONT OF THE MACHINE. PORTABLE: WORK IS CLAMPED IN A FIXTURE, AND A PORTABLE WELDING HEAD IS MOVED OVER THE SEAM. THIS TYPE OF MACHINE IS USED FOR WORKPIECES THAT ARE TOO BULKY TO BE HANDLED BY REGULAR MACHINES

FIG. 8 RELATION OF ELECTRODE WHEELS TO OTHER COMPONENTS IN CIRCULAR AND LONGITUDINAL RSEW MACHINES

Electrodes Electrode wheels made of Resistance Welder Manufacturers Association (RWMA) class 1 material were once used for seam welding of aluminum and magnesium alloys, galvanized steel, and tin-plated steel. However, class 1 material is now uncommon because of the restrictions on cadmium exposure. Typically, class 2, 3, or 20 materials are used (similar to spot welding electrodes). Electrode wheels range in diameter from 50 to 610 mm (2 to 24 in.). Table 4 shows the diameters and body widths of commonly used electrodes, in relation to the size of the welding machine. Narrower wheels are used in machines with knurl or friction drive; wider wheels, in gear drive and idler machines.

TABLE 4 COMMON SIZES OF ELECTRODE WHEELS USED FOR RESISTANCE SEAM WELDING APPLICATIONS

MACHINE SIZE WHEEL DIAMETER WHEEL WIDTH(A) in. mm in. mm 3 SMALL 18 7 9.5 MEDIUM

20

8

9.5-13

LARGE

25-30

10-12

9.5-19

(A)

8 3 1 8 2 3 3 8 4

DATA ARE FOR BODY WIDTH.

A small-diameter electrode wheel (Fig. 9a), or a large-diameter wheel mounted on a canted axis (Fig. 9b), can be used to avoid interference with a sidewall when a narrow flange is being welded on an inside radius or on a re-entrant curve. Beads, fibs, and other extensions on the surfaces to be seam welded can be passed over by cutting notches in the electrode wheels (Fig. 10).

FIG. 9 UPPER ELECTRODE WHEELS DESIGNED TO AVOID INTERFERENCE WITH A SIDEWALL. (A) SMALLDIAMETER WHEEL. (B) CANTED LARGE-DIAMETER WHEEL

FIG. 10 NOTCH ELECTRODE WHEEL USED TO SEAM WELD A WORKPIECE HAVING AN OBSTRUCTION IN THE PATH OF THE WHEEL

Process Variations Seam welding operations vary primarily in terms of travel speed and duration of the welding current. In continuousmotion welding, the electrodes (or workpieces) are driven at a constant speed, and welding current is either interrupted or

flows continuously. In either case, a continuous, leaktight seam is generated. This method is generally used for metals less than about 4.8 mm (

3 in.) thick. 16

In intermittent-motion welding, rotation of the electrode wheels is automatically stopped to produce a spot weld; the electrodes are then rotated to move the workpiece the required distance for the next weld. This method is often used for aluminum and steel alloys less than 4.8 mm (

3 3 in.) thick and selected metals greater than 4.8 mm ( in.) thick. 16 16

Lap Seam Welding. In this process (also known as stitch or roll welding), the on and off cycling of the current while the

electrodes continuously rotate actually determines the weld spacing (Fig. 11a). The minimum joint overlap is the same as for spot welding. There is generally no forging (that is, mash-down) with lap seam welding.

FIG. 11 PROCESS VARIATIONS OF SEAM WELDING. (A) LAP SEAM WELDING. (B) MASH SEAM WELDING. (C) METAL FINISH SEAM WELDING. (D) ELECTRODE WIRE SEAM WELDING. (E) FOIL BUTT SEAM WELDING

Mash seam welding generally has less overlap than lap seam welding and produces more mash-down (electrode force)

with its continuous current (Fig. 11b). Joint thicknesses after welding usually vary from 120 to 150% of the sheet thickness. Continuous seams produced by this method have smooth appearance, but rigid workholding fixtures are required to minimize distortion.

Metal Finish Seam Welding. This process uses a chamfered electrode wheel, and one side of the joint is mashed down while the other side of the joint has a smooth, flat finish (Fig. 11c). Overlap is greater than for mash seam welding, and higher current amperages and electrode forces are also required. Electrode wire seam welding is an internationally patented process that requires an intermediate wire electrode

between each wheel electrode and the workpiece (Fig. 11d). The electrode wire seam welding process is used almost

exclusively for seam welding of tin mill products to fabricate cans. The copper wire travels around the wheel electrodes at the welding speed. The circular or flat cross-section of copper wire electrode provides a continuously renewed surface, but it is not consumed in the welding operation. The tin build-up which would occur on a copper wheel electrode is thus avoided. The process requires specially designed welding systems. The seam welds may be made with two wheel electrodes, or one wheel electrode and a mandrel electrode. Foil Butt Seam Welding. This process requires metal foil to be placed on one or both sides of a butt joint prior to seam welding (Fig. 11e). The foil distributes current to the edges of the sheet, provides additional electrical resistance, helps to contain the molten metal, and also serves as a filler.

Applications Resistance seam welding can be applied to a variety of workpiece shapes. Girth welds can be made in round, square, or rectangular parts by using electrode wheels of suitable diameter. Longitudinal welds can also be made. Advantages of RSEW, as compared with resistance spot and projection welding, are: • •

•

GASTIGHT OR LIQUID-TIGHT JOINTS CAN BE PRODUCED OVERLAP CAN BE LESS THAN FOR RESISTANCE SPOT WELDS OR PROJECTION WELDS, AND SEAM WIDTH CAN BE LESS THAN THE DIAMETER OF SPOT OR PROJECTION WELDS (TABLE 5) ROLLED SECTIONS ARE FREQUENTLY WELDED IN-LINE, THEREBY ELIMINATING SUBSEQUENT OPERATIONS

TABLE 5 RECOMMENDED PRACTICE FOR SEAM WELDING OF 1010 LOW-CARBON STEEL Data are for material free of scale, oxide, paint, grease, and oil. Welding conditions determined by thickness of thinnest outside piece. THICKNESS,T, OF THINNEST OUTSIDE PIECE(A) mm 0.25

in. 0.010

0.53

0.021

0.80

0.031

1.02

0.040

1.27

0.050

1.57

0.062

1.98

0.078

2.39

0.094

2.77

0.109

3.18

0.125

ELECTRODE DIMENSIONS WHEEL FACE WIDTH WIDTH (W)(B) (W)(B) mm in. mm in. 9.5 4.8 3 3 8 16 9.5 4.8 3 3 8 16 13 6.4 1 1 2 4 13 6.4 1 1 2 4 13 7.9 1 5 2 16 13 7.9 1 5 2 16 16 9.5 5 3 8 8 16 11 5 7 8 16 19 13 3 1 4 2 19 13 3 1 4 2

NET ELECTRODE FORCE

HEAT TIME, CYCLES (60 HZ)

COOL TIME (PRESSURE TIGHT), CYCLES (60 HZ)

WELDING SPEED

WELDS PER MM

WELDS PER IN.

APPROXIMATE WELDING CURRENT, A

MINIMUM CONTACTING OVERLAP(C)

kn 1.8

lbf 400

2

1

mm/min 200

in./min 80

0.60

15

8,000

MM 9.5

2.5

550

2

2

190

75

0.48

12

11,000

11

3.1

700

3

2

185

72

0.40

10

13,000

13

4.0

900

3

3

170

67

0.36

9

15,000

13

4.7

1050

4

3

165

65

0.32

8

16,500

14

5.3

1200

4

4

160

63

0.28

7

17,500

16

6.7

1500

6

5

140

55

0.24

6

19,000

17.5

7.6

1700

7

6

125

50

0.24

6

20,000

19

8.7

1950

9

6

120

48

0.20

5

21,000

20.5

9.8

2200

10

7

115

45

0.20

5

22,000

22

IN.

(A) DATA FOR TOTAL THICKNESS OF PILE-UP NOT EXCEEDING 4T. MAXIMUM RATIO BETWEEN THICKNESSES, 3 TO 1. (B) FOR RWMA CLASS 2 ELECTRODE MATERIAL (MINIMUM CONDUCTIVITY, 75% IACS; MAXIMUM HARDNESS, 75 HRB). WHEEL WIDTH IS OVERALL WIDTH OF A FIAT-FACE WHEEL OR OF A WHEEL WITH A 75 MM (3 IN.) RADIUS FACE; FACE WIDTH IS THAT OF THE CONTACTING SURFACE OF FLAT-FACE WHEELS. (C) FOR LARGER ASSEMBLIES, MINIMUM CONTACTING OVERLAP INDICATED SHOULD BE INCREASED 30%.

Limitations of RSEW, apart from those it shares with spot and projection welding, are: • • • •

THE WELD ORDINARILY MUST PROCEED IN A STRAIGHT OR UNIFORMLY CURVED LINE OBSTRUCTIONS ALONG WITH THE PATH OF THE ELECTRODE WHEEL MUST BE AVOIDED OR BE COMPENSATED FOR IN THE DESIGN OF THE WHEEL SHARP CORNER RADII OR ABRUPT CHANGES IN CONTOUR ALONG THE PATH OF THE ELECTRODE WHEEL MUST BE AVOIDED (TABLE 6) FATIGUE LIFE OF RESISTANCE SEAM WELDS IS USUALLY SHORTER THAN THAT OF WELDS MADE BY OTHER SEAM WELDING METHODS

The shorter fatigue life is due to two factors. First, the most common method of cooling the electrode is flood cooling, which quenches the weld. Second, RSEW electrodes move continuously, preventing the application of postweld heat treatment to temper the weld.

TABLE 6 RECOMMENDED MINIMUM RADII OF CURVATURE FOR SEAM WELDED ALUMINUM AND STEEL COMPONENTS

BASE MATERIAL MINIMUM RADIUS OUTSIDE INSIDE RADIUS, R1 RADIUS, R2 MM IN. MM IN. ALUMINUM 250 10 75 3 (A) STEEL 150 6 50 2(A) 57(B) 2 1 (B) 64(C) 4 (D) 1 75 2 (C) 2

3(D) Source: Ref 7

(A) (B) (C) (D)

0.8 MM ≤ T ≤ 1.3 MM (0.031 IN. ≤ T ≤ 0.051 IN.). T ≤ 1.6 MM (T ≤ 0.063 IN.). T ≤ 2.0 MM (T ≤ 0.079 IN.). T ≤ 3.0 MM (T ≤ 0.12 IN.).

Metals Welded. Low-carbon, high-carbon, low-alloy, high-strength low-alloy (HSLA), stainless, and many coated steels

can be resistance seam welded satisfactorily. Alloys with carbon levels above about 0.15% may tend to form areas of hard martensite upon cooling.


7. J.G. BRALLA, ED., HANDBOOK OF PRODUCT DESIGN FOR MANUFACTURING, MCGRAW-HILL, 1986, P 7-39 Procedure Development and Practice Considerations for Resistance Welding

Projection Welding Resistance projection welding (RPW) is a welding process in which current flow and heating are localized at a point or points predetermined by the design or shape of one or both of two parts to be welded (Fig. 12). The process is closely related to resistance spot welding (RSW), in which current flow and heating are localized by one or both electrode contact faces, which determine the location, size, and shape of the weld produced.

FIG. 12 RELATION OF WELD NUGGETS TO EMBOSSED REGIONS OF METAL SHEETS USED IN PROJECTION WELDING. (A) ELECTRICAL CIRCUIT LOOP IS COMPLETED WHEN CURRENT TRAVELS THROUGH EMBOSSED PROJECTION CONTACT POINTS. (B) ADDITION OF PRESSURE AFTER WELDING CURRENT IS TURNED OFF CAUSES PLASTIC DEFORMATION WHICH FLATTENS THE PROJECTIONS.

Projections may be of any practical shape that can properly concentrate the welding current (Fig. 13). In cross-wire projection welding, the curved surfaces of two intersecting wires perform the function of a projection. The shapes of parts may also take the place of conventional projections in projection welding of special types of joints.

FIG. 13 SCHEMATICS SHOWING SELECTED SETUPS WHICH CAN BE USED IN PROJECTION WELDING OPERATIONS

Additional information is available in the article "Projection Welding" in this Volume. Formation of the Weld. The formation of a projection weld nugget, which depends on the design of the projection, the selection of welding conditions, and the adequacy of the resistance welding equipment, is shown in Fig. 14. In this application, 2.33 mm (0.092 in.) thick low-carbon steel is projection welded using embossed spherical projections and a

weld time of 20 cycles (

1 s). 3

FIG. 14 FORMATION OF A WELD NUGGET DURING PROJECTION WELDING OF AN EMBOSSED SPHERICAL PROJECTION. SEE TEXT FOR EXPLANATION.

In the first stage of fabrication (Fig. 14a), the workpieces are brought together under pressure without application of welding current, and the projection may be slightly compressed and indented into the surface of the mating workpiece. Figures 14(b), 14(c), and 14(d) show the stages of formation of a typical weld at 20, 60 to 70, and 100% of the weld time (or heat time). At about 20% of weld time (Fig. 14b), collapse of the projection is nearly complete, and a pressure weld is formed. A nugget of fused weld metal does not begin to form until about 50% of the weld time has elapsed. At about 60 to 70% of weld time (Fig. 14c), fusion has progressed a sufficient distance from the interface to produce a well-defined weld nugget that has about half its final thickness (penetration) and diameter. Sheet separation adjacent to the weld nugget has been

reduced to zero, and the softened metal above the nugget has been flattened against the face of the upper electrode (compare with Fig. 14b). Nugget diameter, penetration, and shear strength continue to increase as weld time progresses. Figure 14(d) shows the fully developed weld nugget. Applicability Projection welding is used where there is a high thickness ratio (for example, nuts welded to sheet), and where the appearance of the sheets must be controlled (for example, appliance surfaces). The principal application of projection welding is the joining of stamped low-carbon, low-alloy, and high-strength low-alloy (HSLA) steel parts. During stamping (punching, drawing, or forming), one of the parts must have a projection which has been formed during the stamping operation to enable projection welding. Projection welding is also used for joining screw-machine parts to stamped parts; the projection is machined or cold formed on the end of the screw-machine part. Projection welding is most successful in workpieces 0.56 to 3.18 mm (0.022 to 0.125 in.) thick. Stock 0.25 mm (0.010 in.) thick has been projection welded; projection design is critical, however, and machines with low-inertia heads and fast follow-up are needed. Sections less than 0.25 mm (0.010 in.) thick are more adaptable to spot welding. Advantages. The principal advantages of projection welding are: •

•

•

• •

•

THE NUMBER OF WELDS THAT CAN BE MADE SIMULTANEOUSLY WITH ONE OPERATION OF THE WELDING MACHINE IS LIMITED ONLY BY THE ABILITY OF THE CONTROLS TO REGULATE CURRENT AND FORCE. BECAUSE OF GREATER CURRENT CONCENTRATION AT THE WELD, AND THUS LESS CHANCE OF SHUNTING, NARROWER FLANGES CAN BE WELDED, AND WELDS CAN BE SPACED CLOSER TOGETHER BY PROJECTION WELDING THAN BY SPOT WELDING. ELECTRODES USED IN PROJECTION WELDING HAVE FACES LARGER THAN THE PROJECTIONS AND LARGER THAN THE FACES OF ELECTRODES USED FOR MAKING SPOT WELDS OF COMPARABLE NUGGET DIAMETER. CONSEQUENTLY, BECAUSE OF LOWER CURRENT DENSITY, ELECTRODES REQUIRE LESS MAINTENANCE THAN DO SPOT WELDING ELECTRODES. PROJECTION WELDS CAN BE MADE IN METAL THAT IS TOO THICK TO BE JOINED BY RSW. FLEXIBILITY IN SELECTION OF PROJECTION SIZE AND LOCATION ALLOWS WELDING OF WORKPIECES IN THICKNESS RATIOS OF 6 (OR MORE) TO 1. WORK-PIECES IN THICKNESS RATIOS GREATER THAN ABOUT 3 TO 1 SOMETIMES ARE DIFFICULT TO SPOT WELD. THE PROCESS CAN BE USED TO MAKE LEAKPROOF JOINTS (FOR EXAMPLE, RING PROJECTIONS).

Limitations of projection welding include: • •

•

•

FORMING OF ONE OR MORE PROJECTIONS ON ONE OF THE WORKPIECES MAY REQUIRE EXTRA OPERATIONS UNLESS THE PARTS ARE PRESS FORMED TO DESIGN SHAPE. WHEN SEVERAL WELDS ARE MADE AT ONCE WITH THE SAME ELECTRODE, ALIGNMENT OF THE WORK AND DIMENSIONS (PARTICULARLY HEIGHT) OF THE PROJECTIONS MUST BE HELD TO CLOSE TOLERANCES TO OBTAIN CONSISTENT WELD QUALITY. WHEN MAKING SIMULTANEOUS PROJECTION WELDS, THE PROJECTION LAYOUT WILL BE DICTATED BY THE SHUNT CURRENT PATHS, WHICH MAY NOT COINCIDE WITH THE DESIRED LOCATION. BECAUSE TWO PARTS ARE MOVING RELATIVE TO EACH OTHER DURING THE WELDING PROCESS, PROPER GAGING IS DIFFICULT.

Cross-Wire Welding

Resistance welding of crossed wires is a type of projection welding process. In practice, it usually consists of welding a number of parallel wires at right angles to one or more other wires or rods. There are many specific ways to perform the welding operation, depending upon production requirements, but the final finished product is essentially the same regardless of the method used. Typical crossed-wire products include items such as stove and refrigerator racks, grills of all kinds, lamp shade frames, poultry equipment, wire baskets, fencing, grating and concrete rein-forcing mesh. Welding Machines Press-type machines, with either single-phase or three-phase transformers, usually are used for projection welding. The welding head in these machines is guided by bearings or ways and moves in a straight line. Platens with T-slots or tapped holes are used for mounting the welding dies or electrodes. Rocker-arm machines generally are not used for projection welding because the electrode moves in an arc that can cause slippage between the components as the projection collapses. At the start of the welding stroke (Fig. 15a), equal air pressure is applied to the top of the piston and to the diaphragm. By the time the piston has completed its travel, the diaphragm has been compressed by the retraction of the internal shaft. Retraction of the internal shaft simultaneously compresses the diaphragm and closes the welding-current switch; the control lever for the switch rides in the shaft collar.

FIG. 15 SCHEMATIC SHOWING MOTION OF INTERNAL COMPONENTS IN A LOW-INERTIA WELDING HEAD USED FOR PROJECTION WELDING OPERATIONS

When the welding current is initiated (Fig. 15b), the only remaining mass to be moved is the internal shaft and its attached electrode. Air pressure acting on the diaphragm and the force of the compressed spring between the inner and outer shafts easily overcome the low inertia of the system and move the upper electrode downward as the projection collapses. Thus, the workpieces are kept in close contact under partial electrode force until the welding head follows and forges the weld region with the full electrode force (Fig. 15c). This setup, however, is no longer used because diaphragm cylinders are now classified as obsolete technology. Process Variables The major variables that affect projection welding are welding current, electrode force, and weld time. Specifications for projection welding of stainless steels relative to low-carbon steel from 0.36 to 3.18 mm (0.014 to 0.125 in.) thick are given in Table 7. These data are intended to serve as starting points and should be adjusted to suit the specific application and the equipment used.

TABLE 7 RECOMMENDED ELECTRODE FORCE AND WELDING CURRENT SPECIFICATIONS FOR PROJECTION WELDING OF SELECTED STAINLESS STEELS RELATIVE TO 1010 LOW-CARBON STEEL

Parameters for projection welding of low-carbon steel wire are given in Table 8.

TABLE 8 GUIDELINES FOR PROJECTION WELDING OF LOW-CARBON STEEL WIRE WIRE DIAMETER

CONDITIONS FOR 15% SETDOWN(A) ELECTRODE WELDING WELD CURRENT, FORCE STRENGTH A KN LBF KN LBF

CONDITIONS FOR 30% SETDOWN(A) ELECTRODE WELDING WELD CURRENT, FORCE STRENGTH A KN LBF KN LBF

CONDITIONS FOR 50% SETDOWN(A) WELD ELECTRODE WELDING STRENGTH CURRENT, FORCE A KN LBF KN LBF

0.45

100

600

2.0

450

0.67

150

800

2.2

500

0.89

200

1,000

2.4

550

0.56

125

1,800

4.3

975

1.2

260

2,650

5.0

1,125

1.6

350

3,400

5.6

1,250

1.6

360

3,300

8.9

2,000

2.7

600

5,000

10.7

2,400

3.3

750

6,000

11.1

2,500

2.6

580

4,500

16.4

3,700

3.8

850

6,700

18.7

4,200

5.5

1240

8,600

19.6

4,400

3.7

825

6,200

22.7

5,100

6.4

1450

9,300

27.1

6,100

8.9

2000

11,400

28.9

6,500

4.9

1100

7,400

29.8

6,700

9.2

2060

11,300

37.1

8,350

13.3

3000

14,400

39.1

8,800

6.2

1400

9,300

42.7

9,600

12.9

2900

13,800

50.3

11,300

19.8

4450

17,400

52.9

11,900

7.6

1700

10,300

54.3

12,200

15.1

3400

15,800

60.5

13,600

23.6

5300

21,000

64.9

14,600

0.45

100

600

1.6

350

0.67

150

800

1.8

400

0.89

200

1,000

2.0

450

10

0.56

125

1,850

3.3

750

1.2

260

2,770

3.8

850

1.6

350

3,500

4.0

900

17

1.6

360

3,500

6.7

1,500

2.7

600

5,100

7.6

1,700

3.3

750

6,300

8.0

1,800

23

2.6

580

4,900

12.5

2,800

3.8

850

7,10

13.3

3,000

5.5

1240

9,000

13.8

3,100

30

3.7

825

6,600

20.5

4,600

6.4

1450

9,600

22.2

5,000

8.9

2000

12,000

23.6

5,300

40

4.9

1100

7,700

27.6

6,200

9.2

2060

11,800

30.2

6,800

13.3

3000

14,900

32.0

7,200

WELD TIME(B), CYCLES

MM IN. COLD-DRAWN WIRE 1.6 5 1 16 10 1 8 4.8 17 3 16 6.4 23 1 4 7.9 30 5 16 9.5 40 3 8 11 50 7 16 13 60 1 2 HOT-DRAWN WIRE 1.6 5 1 3.2

3.2 4.8 6.4 7.9 9.5

16 1 8 3 16 1 4 5 16 3 8

11 13

7 16 1 2

50

6.2

1400

10,000

39.1

8,800

12.9

2900

14,000

42.7

9,600

19.8

4450

18,000

45.4

10,200

60

7.6

1700

11,000

51.1

11,500

15.1

3400

16,500

55.2

12,400

23.6

5300

22,000

57.8

13,000

Source: Ref 9 (A) (B)

SETDOWN, % = (DECREASE IN JOINT HEIGHT ÷ DIAMETER OF SMALLER WIRE) × 100. FOR 15.30, AND 50% SETDOWN.

Welding current required for projection welding, although slightly less per weld than that needed for spot welding, must

be high enough to cause fusion before the projection is completely flattened. The recommended current is the highest current that, when used with the current electrode pressure, does not cause excessive expulsion of metal. Electrode force used in projection welding depends on the work metal, the size and design of the projection, the number

of projections in the joint, and the welding machine. Excessive force causes the projections to collapse before the weld area has reached the proper temperature, resulting in the formation of ring welds, in which fusion occurs around the periphery of the projection but is incomplete at the center. For best weld appearance, the electrode force should be such that the projection is flattened completely after the metal has reached welding temperature. Weld time for a given type and thickness of work metal depends on welding current and rigidity of the projection. Weld

time is less important than electrode force in projection welding of low-carbon, low-alloy, and HSLA steel, provided that the time is sufficient to produce a nugget of adequate size at the chosen welding current. A short weld time creates higher production efficiency and less discoloration and distortion of the workpiece. After the proper electrode force and welding current are determined, the weld time is adjusted to make the desired weld. Electrodes An electrode designed for RSW can be used for projection welding if the electrode face is large enough to cover the projection being welded, or the pattern of projections being welded simultaneously. To minimize marking and indentation of workpieces, the recommended electrode-face diameter for making a single projection weld usually is two or more times the diameter of the projection. In multiple-projection welding, the electrode face should be large enough to extend beyond the boundaries of the pattern of projections by approximately the diameter of one projection. RWMA class 2 electrode materials generally are preferred as general-purpose materials because they provide the best compromise among electrical and conductivity, strength, hardness, and temperature resistance. Typically, class 10, 11, or 12 materials are used on the bottom side of the weld to resist damage from the high localized force.


9. RESISTANCE WELDING MANUAL, 3RD ED., VOL 1, RESISTANCE WELDER MANUFACTURERS ASSOCIATION, 1956, P 88 Procedure Development and Practice Considerations for Resistance Welding

Flash Welding Flash welding (FW) commonly is used to join sections of metals and alloys in production quantities. It is a resistance/forge welding process in which the items to be welded are securely clamped to electric current-carrying dies, heated by the electric current, and upset (Fig. 16). Clamping ensures good electrical contact between the current-carrying dies and the workpiece and prevents the parts from slipping during the upsetting action. Flash welding equipment must be durable to withstand high clamping force and upset pressures without deflecting. If deflection occurs, misalignment of workpieces may occur during welding. Flash welding is rapid and economical, and when properly executed, welds of uniform high quality are produced. A typical flash welding machine comprises a horizontal press-transformer combination with conducting work-clamping dies mounted on the press platens.

FIG. 16 WELD PRODUCED WHEN USING THE FLASH WELDING PROCESS. (A) WORKPIECES SECURELY CLAMPED IN CURRENT-CARRYING DIES BEFORE UPSETTING OPERATION IS INITIATED. (B) FINISHED WELD PRODUCED AFTER UPSETTING OPERATION

Additional information is available in the article "Flash Welding" in this Volume. Applications Flash welding can be used for joining many ferrous and nonferrous alloys and combinations of dissimilar metals. In addition to low-carbon steels, metals that are flash welded on a production basis include low-alloy steels, tool steels, stainless steels, aluminum alloys, magnesium alloys, nickel alloys, and copper alloys. Table 9 gives typical upsetting pressures for numerous classes of alloys.

TABLE 9 RECOMMENDED UPSETTING PRESSURES FOR FLASH WELDING OF SELECTED ALLOYS

LOW FORGING

ALLOYS UPSET PRESSURE MPA KSI SAE DESIGNATION 69 10 1020, 1112, 1315

MEDIUM FORGING

103

15

HIGH FORGING

172

25

STRENGTH CLASSIFICATION

1045, 1065, 1335, 3135, 4130, 4140, 8620, 8630 4340, 4640, 300M

OTHERS HIGH-STRENGTH LOWALLOY (HSLA) STEELS

TOOL STEELS, TITANIUM, ALUMINUM

EXTRA-HIGH FORGING

241

35

A-286, 19-9 DL

NICKEL-BASE ALLOYS, (A)

Source: Ref 10

(A) MATERIALS EXHIBITING EXTRA-HIGH COMPRESSIVE STRENGTH AT ELEVATED TEMPERATURES. Welding Parameters Fundamentally, flash welding involves heating the ends of the pieces to be welded and subsequently forging them together. During the heating phase, a thermal distribution pattern is established along the axial length of the pieces being joined, which is characterized by a steep temperature gradient. The major difference between the temperature pattern developed in flash welding and that developed in resistance welding is that flash welding produces a much steeper thermal gradient. This steep thermal gradient, combined with the resulting characteristic upset pattern, enables flash welding to accommodate a much greater variety of materials and shapes than can be welded by resistance welding. Once the proper temperature-distribution pattern has been established, the abutting surfaces are rapidly forced together. The parts must be securely held together during the forging process to prevent slipping. Three distinct peaks are characteristic of flash welds (Fig. 17). The two peaks on either side of the weld line represent the material displaced by the upsetting action; the center peak is the molten metal extruded out of the weld, including oxides or contaminants formed during heating.

FIG. 17 SCHEMATIC SHOWING CROSS-SECTIONAL VIEW OF TYPICAL PEAKS AND FLOW LINES GENERATED IN A FLASH WELD

Heat (Energy) Sources One of the major considerations for flash welding is the electrical power service. Flashing is a term used to describe the major heating process during flash welding. When the ends of the workpiece are brought together under light pressure, an electrical short circuit is established through the material. Because the abutting surfaces are not perfectly matched, the short-circuit current flows across the joint only at a few small contact areas. The large amount of current flowing through a relatively small area causes very rapid heating to the melting point. Heating is so rapid and intense that molten metal is expelled explosively from the joint area. Following this expulsion, a brief period of arcing occurs. Arcing is not sustained due to the low voltages employed. Studies have shown that stable flashing can be maintained with as low as 2 but typically the flashing voltage for alternating-current machines is from 5 to 10 V.

1 to 3 V, 2

Following the expulsion of molten metal and subsequent arcing, small craters are formed on the ends of the abutting surfaces. The pieces are steadily advanced toward one another, and other short circuits are formed and additional molten metal is expelled. This process continues as random melting, arcing, and expulsion occur over the entire cross-sectional surface. During flashing, many active areas are in various stages of this sequence (Fig. 18). Flashing surfaces act as heat sources, and the steep thermal profile is established primarily from these heat sources. Temperatures of the flashes are at or above the melting point of the material and are progressively lower as distance progresses from the flashing surface toward the clamp.

FIG. 18 GROWTH OF CONTACT POINTS PRODUCED BY ENERGY INPUT DURING FLASH WELDING

Force Parts to be welded must be clamped or fixtured securely to provide good electrical contact with current-carrying dies and to transmit the upset forces. Also, a reliable source of force for the upsetting action is required. Systems for generating these forces vary greatly and are determined by the cross sections to be welded. The simplest welding machines derive their power for clamping and upsetting from the operator (that is, the parts are clamped to the dies via screw, lever, cam, or toggle force-multiplying linkages). Upset forces are generated in a similar manner. In larger machines, these forces are generated by pneumatic systems, oil hydraulic systems, combination air/oil hydraulic systems, or motor-driven cam systems. Machines Machines used for flash welding generally consist of a mainframe; a low-impedance welding transformer; a stationary platen; a movable platen on which clamping dies, electrodes, and other tools needed to position and hold the work-pieces are mounted; flashing and upsetting mechanisms; and the necessary electrical, air, or hydraulic controls (Fig. 19).

FIG. 19 FLOW DIAGRAM SHOWING ADAPTIVE CONTROLS USED TO CONTROL AND MONITOR FLASH WELDING PROCESS

Flash welding machines may be manual, semiautomatic, or fully automatic. With manual operation, the operator controls the speed of the platen from the time flashing is initiated until the upset is completed. In semiautomatic operation, the operator manually initiates flashing and then actuates an automatic cycle that completes the weld. In fully automatic operation, after the operator initiates the welding sequence, a fully automatic cycle can be used through the use of position-indicating devices and timers. Values for the various welding parameters are preselected by the operator. Automatic feedback control is used in some applications. These fully adaptive circuits vary from current to voltage feedback circuits, and data are obtained to control the welding sequences. Figure 19 shows various adaptive controls used by different flash welding unit manufacturers. Transformers are used to supply electrical energy for the welding operation. The transformer tap switch frequently is a rotary, eight-step, knife-type, fully enclosed locking switch, which provides convenient adjustment of the welding voltage to suit work requirements. Voltages from 4 to 16 V are common.

Clamping Dies and Fixtures Workpieces must be accurately clamped to maintain alignment, to allow the secondary current to pass into the workpieces, and to apply the upsetting force properly. Generally, the parts of the clamping mechanism that actually grip the workpiece are the electrodes, often called clamping dies. The surface of the material in contact with the dies must be cleaned of scale, rust, paint, or contaminants to ensure proper current flow. The materials commonly used for clamping dies are RWMA class 3 and hardened tool steels such as H11, L6, and O1. Bronze and other copper-base electrode materials can be used in some applications. The die half (upper or lower) that

conducts the current usually is made of a copper-base material. The other half is made of the same copper-base material or of hardened steel. Cooling of Electrodes. The need for water cooling of flash welding electrodes depends on the size of the electrode, die,

or fixture, the magnitude of the flashing current, and the production rate. When the mass of the electrodes is large compared with that of the workpiece, the heat sink generally is large enough to dissipate the heat generated by the resistance to current flow and the heat absorbed from the workpiece. With high flashing currents and high production rates, electrodes generally must be cooled. Shape and Size of Clamping Dies. Generally, the shape of the clamping-die surface is such that the die encompasses almost the entire workpiece surface. The required area of clamping-die contact depends on the current needed for heating the workpiece and the pressure needed for holding it. Semicircular dies are used where the line contact provided by Vshape dies gives insufficient surface for the current to flow without burning the workpiece, or for holding the workpiece without marking it.

Table 10 gives the minimum lengths of clamping dies for welding rounds, tubing, and other sections of various diameters or minimum dimensions; these data are for welding steels of low or medium forging strength.

TABLE 10 MINIMUM LENGTHS OF CLAMPING DIES, WITH AND WITHOUT BACKUP, FOR FLASH WELDING SELECTED DIAMETERS OF WORKPIECES MADE FROM STEELS OF LOW OR MEDIUM FORGING STRENGTH WORKPIECE DIAMETER(A)

in. 2.00

MINIMUM LENGTH OF CLAMPING DIE WITH BACKUP mm in. 31.8 1.25

mm 152

in. 6.00

MINIMUM LENGTH OF CLAMPING DIE WITH BACKUP in. mm 82.3 3.25

2.50 3.00 3.50 4.00 4.50 5.00 5.50

44.5 50.8 57.2 63.5 69.9 69.9 76.2

165 178 191 203 216 229 241

6.50 7.00 7.50 8.00 8.50 9.00 9.50

88.9 95.3 102 108 114 121 127

WORKPIECE DIAMETER(A)

MINIMUM LENGTH OF CLAMPING DIE WITH WITHOUT BACKUP BACKUP(B)

WORKPIECE DIAMETER(A)

mm 6.35, 7.92 9.53 12.70 19.05 25.40 38.10

mm 9.53

in. 0.375

mm 25.4

in. 1.00

mm 50.8

9.53 9.53 12.70 19.05 25.40

0.375 0.375 0.500 0.750 1.000

38.1 44.5 50.8 63.5 76.2

1.50 1.75 2.00 2.50 3.00

63.5 76.2 88.9 102 114 127 140

in. 0.250, 0.312 0.375 0.500 0.750 1.000 1.50

1.75 2.00 2.25 2.50 2.75 2.75 3.00

3.50 3.75 4.00 4.25 4.50 4.75 5.00

(A) DIAMETER OF ROUNDS OR TUBING, OR MINIMUM DIMENSION OF OTHER SECTIONS. (B) BACKUP IS RECOMMENDED FOR ALL WORKPIECE DIAMETERS OR MINIMUM DIMENSIONS OVER 38.1 MM (1.50 IN.).

Preweld Processing Parts that are flash welded come in a variety of forms, including forgings; rolled or extruded bar; sheet, strip, or plate, ring preforms; and castings. Each of the above parts or assemblies requires preweld preparation. Cleaning. As a minimum precaution, the ends of the workpiece that are clamped on the current-conducting die must be free from dirt, scale, surface oxidation, and grease. In addition, the ends of the workpiece that extend into the weld zone must be free from any contamination that could react with the base metal at the high temperatures developed during welding. Cleaning is needed because of the high current densities developed in the workpiece at the current-conducting dies. If insufficient electrical contact is made, weld quality is poor and localized hot spots can develop between the current-conducting dies and the workpiece. These localized hot spots are called "die burns." Many of the alloys welded have tightly adherent, highly resistive oxides that must be removed prior to welding.

In addition, poor fit-up, loose scale, and grease may cause the parts to slip during upsetting. Common cleaning techniques include (1) abrading the surfaces by grinding, grit blasting, or wire brushing; (2) pickling or chemical descaling; and (3) vapor degreasing.

End Preparation. Die burns, upset slippage, and inferior welds may be caused by poor fit-up between the current-

conducting dies and the workpiece. In some cases, the ends of the workpiece must be machined or ground to fit the dies, particularly in welding of rough forgings or castings. Also, a chamfer may be required on the ends of the workpiece to initiate flashing action. Alignment of Workpiece. After the workpieces are clamped in the welding unit, the alignment of the workpieces with

respect to each other along the axis of upset must be maintained. After the flashing process and during upsetting, any misalignment may cause the parts to overlap one another or cause a skewed weld. If this occurs, insufficient metal will be extruded out of the weld zone during upsetting and a poor-quality weld will be formed (Fig. 20).

FIG. 20 EFFECT OF WORKPIECE ALIGNMENT ON JOINT QUALITY AND WELD UPSET

Dimensional Tolerances. Some combinations of joint length and stock thickness are difficult or even impossible to

weld. Table 11 gives relationships between joint length (sheet width) and sheet thickness that have proven successful in a large number of production jobs. Relationships of tube diameter to wall thickness that have been successfully welded are listed in Table 12.

TABLE 11 LENGTHS OF JOINTS COMMONLY FLASH WELDED IN SELECTED THICKNESSES OF FLAT SHEET

SHEET THICKNESS mm in. 0.25 0.010 0.51 0.020 0.76 0.030 1.02 0.040 1.27(A) 0.050(A) 1.52(A) 0.060(A) (A) 2.03 0.080(A) 2.54(A) 0.100(A) 3.18(A) 0.125(A) (A) 4.75 0.187(A)

JOINT LENGTH in. mm 25 1 125 5 255 10 380 15 510 20 635 25 890 35 1145 45 1458 57 2235 88

(A) LENGTH OF JOINTS IN STOCK THICKNESSES OF 1.27 MM (0.050 IN.) AND GREATER CAN BE UP TO 2540 MM (100 IN.) DEPENDING ON MATERIAL EXTENSION, PLATEN TRAVEL, DIE ALIGNMENT, AND CLAMPING. TABLE 12 MAXIMUM FLASH-WELDABLE DIAMETERS OF TUBING OF SELECTED WALL THICKNESSES

WALL THICKNESS MAXIMUM DIAMETER in. mm in. mm 1 0.51 0.020 13 0.76

0.030

19

1.27

0.050

32

1.57

0.062

38

2.03 2.54 3.18 4.75 6.35 9.52 12.7

0.080 0.100 0.125 0.187 0.250 0.375 0.500

50 75 100 150 230 380 510

2 1 4 1 4 1 1 2

1

2 3 4 6 9 15 20

Burnoff. After the parts are clamped in the welding machine and the welding power is activated, the abutting surfaces are brought together for a brief period of violent flashing. This phase of the welding sequence is called "burnoff." It serves the function of squaring off the abutting surfaces and compensates for inconsistencies in end preparation. Preheating is a resistive heating phase of the welding process in which the heat is generated by the electrical resistance

of the workpieces. It is accomplished by bringing the workpieces together under light pressure, creating a short circuit. The pressure must be great enough to prevent flashing, but not so great that workpieces are prematurely welded. In practice, preheating is performed in a cyclic manner, workpieces are brought together briefly, then separated for a brief period to allow the heat generated to diffuse into them. This sequence is repeated several times. Fully automatic machines perform this function, thus eliminating operator variables. Flashing. The primary purpose of flashing is to generate enough heat to produce a plastic zone that permits adequate

upsetting. The rate of energy input must be in proper proportion to the travel of the platen or movable die, so that constant flashing is maintained until the appropriate amount of metal is flashed off and the required plastic zone is obtained.

Upsetting (Forging). Bonding takes place during the upsetting action, and some metal must be extruded from the weld

zone to remove slag and other inclusions not expelled during flashing. The extruded metal must extend beyond the crosssectional boundaries of the workpiece to ensure that maximum amounts of slag and inclusions are removed when the weld upset is removed during subsequent trimming. When the weld upset is removed, no evidence of the weld should remain. Porosity near the outer surface on an etched section of the weld, or a crevice around the workpiece after the weld upset has been removed, indicates incomplete bonding because of either insufficient upsetting force or insufficient plasticity during upsetting. A large upset produced by an excessive upsetting force can be detrimental to weld quality, can waste metal, and indicates that most of the plastic metal has been extruded, requiring bonding to take place in metal where plasticity may not have been sufficient to ensure a good weld.


10. R.L. O'BRIEN, ED., WELDING HANDBOOK, 8TH ED., VOL 2, AWS, 1991, P 593 Procedure Development and Practice Considerations for Resistance Welding

Upset Welding (Ref 11) Upset welding (UW) is a resistance welding process that produces coalescence over the entire area of faying surfaces or progressively along a butt joint by the heat obtained from the resistance to the flow of welding current through the area where those surfaces are in contact. Pressure is applied before heating is started and is maintained throughout the heating period (Fig. 21). The equipment used for upset welding is very similar to that used for flash welding. It can be used only if the parts to be welded are equal in cross-sectional area. Upset welding differs from flash welding in that the parts are clamped in the welding machine and force is applied to bring the parts tightly together. High-amperage current is then passed through the joint, which heats the abutting surfaces. When they have been heated to a suitable forging temperature, an upsetting force is applied and the current is stopped. The high temperature of the work at the abutting surfaces, plus the high pressure, causes coalescence to take place. After cooling, the force is released and the weld is completed. There is no arc or flash in upset welding. The area at the joint is usually enlarged over its original dimension.

FIG. 21 WELD OBTAINED WHEN USING THE UPSET WELDING PROCESS. (A) WORKPIECES SECURELY CLAMPED IN CURRENT-CARRYING DIES WHILE PRESSURE AND CURRENT ARE APPLIED. (B) FINISHED WELD PRODUCED AFTER WELDING OPERATION IS COMPLETED

This process is used for welding small wires, tubing, piping, rings, and strips where the cross-sectional areas of both pieces are identical. If intimate contact is not obtained because of improper joint preparation, the weld will be defective. Additional information is available in the article "Upset Welding" in this Volume.


11. H.B. CARY, MODERN WELDING TECHNOLOGY, 2ND ED., PRENTICE-HALL, 1989, P 251-252 Procedure Development and Practice Considerations for Resistance Welding

High-Frequency Resistance Welding High-frequency resistance welding (HFRW) is a resistance welding process that produces coalescence of metals by the heat generated from the resistance of the workpieces to a high-frequency current (10 to 500 kHz) and the rapid application of an upsetting force once heating is completed. The current is applied to the base metal by using electrodes that are held against the surface of the workpiece. The principal applications of HFRW are for making such structural shapes as pipe and tubing from continuous strips of metal. Two process variations of HFRW are continuous seam welding (Fig. 22a, b, c, d, e, and f) and finite-length butt welding (Fig. 22g), which is a variation of induction upset welding (UW-I).

FIG. 22 TYPICAL HIGH-FREQUENCY WELDING APPLICATIONS USED TO JOIN PLATES AND TO FORM TUBES

Detailed information is available in the article "High-Frequency Welding" in this Volume. Continuous seam welding can be described by its most common application, high-frequency upset welding (UW-HF) of a tube (Fig. 23). A pair of electrodes is placed on either side of an open V formed by the converging edges of the strip being made into tubing. The angle of this V typically varies from 4 to 7°. The edges of the strip converge between the weld pinch rolls, where a certain degree of forging upset is accomplished.

FIG. 23 HIGH-FREQUENCY LONGITUDINAL BUTT SEAM WELDING OF A TUBE

As the material moves through the rolls from the contacts to the point of welding, high-frequency current introduced at the electrodes flows on the opposing faying surfaces of the strip, down one side of the V to the point of welding, and back to the other electrode. Because this process is performed at frequencies of about 400 kHz, the depth of heating is quite shallow and heating occurs rapidly. Power is adjusted to bring the edges to a welding temperature at the time they join, and contaminants on the surfaces, as well as most of the molten material, are squeezed out of the seam toward the inside and outside of the tube when the weld is completed. This upset material can be trimmed off flush to the surfaces of the tube if desired. The other applications of continuous high-frequency resistance seam welding (Fig. 23) vary from tube welding only in the shape of the workpiece. Contacts introduce current at the open part of a V that is being closed by rolls at the point at which the material has been heated to a high enough temperature to weld. Finite-Length Welding. The process shown in Fig. 22(g), in which a butt joint of finite length is welded between two

parts, is used for applications in coil end joining as well as in the manufacture of strip and sheet blanks for subsequent press forming. In this process, heating is accomplished along a seam that has already been closed under light pressure. The path of the heating current is controlled by the use of a proximity conductor positioned over the seam. The highfrequency heating causes the edges to be brought to welding temperature, after which pressure is increased to weld the joint. The appropriate frequency is selected to achieve heating through the material thickness. For example, 10 kHz can be used up to 6.4 mm (0.25 in.) thickness. Joints made this way can be produced very quickly, typically in 1 s. Compared to flash welding, the amount of upset is small, no shower of sparks and debris is emitted from the weld, and very little metal is consumed in making the joint. Welding Parameters. Table 13 gives typical HFRW weld speeds as a function of wall thickness for tube and pipe

products.

TABLE 13 TYPICAL WELD RATES FOR HIGH-FREQUENCY RESISTANCE WELDING OF TUBE AND PIPE

SPEED WALL THICKNESS 200 KW, KW, 600 76 MM (3.0 IN.) 305 MM (12 IN.)

MM 1.52 2.03 3.05 4.06 6.10 8.13 9.91 11.94 13.97 16.00

IN. 0.06 0.08 0.12 0.16 0.24 0.32 0.39 0.47 0.55 0.63

DIAMETER M/S FT/MIN 2.91 571 2.17 428 1.45 285 1.08 213 0.67 132 0.41 81 ... ... ... ... ... ... ... ...

DIAMETER M/S FT/MIN ... ... ... ... ... ... ... ... 2.00 393 1.38 272 0.97 191 0.70 137 0.50 99 0.35 69

Procedure Development and Practice Considerations for Resistance Welding

Capacitor Discharge Stud Welding The capacitor discharge process uses direct current produced by the rapid discharge of a capacitor bank to create an electric arc and the heat needed for melting both stud and workpiece. The weld time is very short, ranging from 1 to 6 ms. Because the weld time is so short, no ferrule or fluxing is required. The equipment consists of a stud gun, capacitor discharge power/control unit, and suitable weld cable. Figure 24 illustrates a portable capacitor discharge stud welding system.

FIG. 24 KEY COMPONENTS OF A PORTABLE CAPACITOR DISCHARGE STUD WELDING SYSTEM

The three different methods of capacitor discharge stud welding include: initial contact, initial gap (Fig. 25), and drawn arc. The difference among these three methods is in the manner of arc initiation. The initial contact and initial gap methods utilize a small, specially engineered projection or tip to obtain arc initiation, whereas the drawn arc method initiates the arc by drawing the stud away from the workpiece.

FIG. 25 SEQUENCE OF OPERATIONS REQUIRED TO WELD A STUD TO A WORKPIECE USING THE INITIAL GAP METHOD OF CAPACITOR DISCHARGE STUD WELDING

Studs. Low-carbon steel, stainless steel, aluminum, and brass are the materials used most often for studs welded

commercially with the capacitor discharge method. The preferred weld base of a capacitor discharge welding stud is round, and the shank may be of almost any shape--threaded, unthreaded, round, square, rectangular, tapered, or grooved. The maximum weld-base diameter being welded commercially with the capacitor discharge welding process is 13 mm (

1 2

in.). Weld Current/Weld Time Relationships. The weld current process ranges from approximately 600 to 25,000 A,

depending on stud size and weld method. The weld current range with the drawn arc method is approximately 600 to 3000 A, whereas with the initial contact or gap method the range is approximately 2500 to 25,000 A. Weld time with the drawn arc method is generally in the range of 4 to 6 ms, whereas with the initial contact and gap method it is generally in the range of 1 to 3 ms. Metals Welded. The capacitor discharge stud welding process is an arc welding process, but differs from other arc

welding processes in that the weld is made in a few milliseconds. For this reason, acceptable metallurgical results can be obtained without shielding the arc or using a flux. The welding of aluminum using the drawn arc method of capacitor discharge stud welding does require a shielding gas, however. Because of the speed at which the weld takes place, very little metal is melted. Weld penetration is slight, and very little stud metal intermixes with the plate metal. Therefore, studs may be welded to very thin metals without melt-through, and many dissimilar metals can be welded with acceptable results. Various stud and plate material combinations can be welded (Table 14).

TABLE 14 BASE METAL/STUD METAL COMBINATIONS THAT CAN BE SUCCESSFULLY JOINED USING CAPACITOR DISCHARGE STUD WELDING

BASE METAL

STUD METAL STAINLESS LOWCARBON STEEL STEEL(A) (300 SERIES)(B)

LEADFREE COPPER ALLOY

ALUMINUM ALLOY (1100 SERIES, 5000 SERIES, 6000 SERIES)

TITANIUM OR TITANIUM ALLOY

LOW-CARBON STEEL (1006-1025) STAINLESS STEEL (300 SERIES, 400 SERIES) ALUMINUM ALLOY (1100 SERIES, 3000 SERIES, 5000 SERIES, 6000 SERIES) COPPER ALLOY (LEADFREE BRASS) ZINC ALLOY (DIE CAST) TITANIUM OR TITANIUM ALLOY

X

X

X

...

...

X

X

X

...

...

...

...

...

X

...

X

X

X

...

...

...

...

...

X

...

...

...

...

...

X

(A) COMPOSITION: 0.23% C (MAX), 0.60% MN (MAX), 0.04% P (MAX), 0.05% S (MAX). (B) WITH EXCEPTION OF FREE-MACHINING TYPE 303 STAINLESS STEEL Procedure Development and Practice Considerations for Resistance Welding

Resistance Welding of Stainless Steels Stainless steels are readily resistance welded by spot, seam, projection, flash, upset, high-frequency, and capacitor discharge. Generally, the weld time and welding current are less than those used for welding carbon steel, but the electrode force is usually greater. Austenitic stainless steels of the 300 series are resistance welded more often than any other metal except lowcarbon steel. Ideally, stainless steel to be resistance welded should contain a maximum of 0.08% C (as in types 304, 316, and 347), although steels of higher carbon content (such as types 301, 302, 309, and 310) can be successfully resistance welded. The martensitic and ferritic types of stainless steel can be welded satisfactorily. The martensitic types are less frequently resistance welded, because joints made in them are hard and brittle in the as-welded condition. Equipment Resistance welding of stainless steel requires less transformer capacity than for equivalent thicknesses of other metal, under the same conditions. The welding machines use a singlephase alternating-current power supply, a three-phase rectifier, or a frequency-converter power supply. A weld made with a three-phase power supply usually requires slightly higher current and a longer weld time than a weld made with a single-phase supply. Factors Affecting Resistance Welding of Stainless Steels The characteristics of stainless steel that affect resistance welding include electrical resistivity, thermal conductivity, melting temperature, strength at elevated temperatures, coefficient of thermal expansion, and contact resistance. Electrical Resistivity. Stainless steels have much higher electrical resistance than carbon steels. As a result, more heat is

generated in a stainless steel with the same current. Therefore, resistance welding of a stainless steel requires lower currents or shorter weld times than welding of a carbon steel. Thermal Conductivity. Stainless steel has a lower thermal conductivity than carbon steel and, therefore, heat is

conducted away from the weld zone more slowly. This, in conjunction with the electrical resistivity, means that less heat has to be applied to reach the melting temperature. Melting temperatures of stainless steels have an effect on the amount of heat required to produce fusion for welding. The austenitic stainless steels melt in various ranges between 1370 and 1455 °C (2500 and 2650 °F), and the martensitic and ferritic alloys melt in ranges between 1400 and 1530 °C (2550 and 2790 °F). Plain low-carbon steels melt at temperatures between 1480 and 1540 °C (2700 and 2800 °F).

Coefficient of Thermal Expansion. Austenitic and nitrogen-strengthened austenitic stainless steel expands and

contracts with changing temperature almost 50% more than does plain carbon steel. These dimensional changes and the slower heat diffusion in austenitic grades result in greater thermal stress, which leads to warping. The ferritic, martensitic, and precipitation-hardening grades of stainless steel have coefficients of thermal expansion from 6 to 11% lower than that of plain carbon steel. High strength at room and elevated temperatures of nitrogen-strengthened stainless steels and, to a lesser extent, of

straight-chromium grades makes it necessary to use greater electrode force than is required for carbon steel to bring the work-metal surfaces together for the required intimate contact at points of welding. Contact resistance of stainless steel is higher than that of carbon steel and, therefore, greater electrode pressure is

needed to make good resistance welds. In addition, the welding process is more stable because the base material is not as susceptible to surface contamination. Welding Characteristics of Stainless Steels Austenitic Stainless Steels. All of the austenitic stainless steels in the solution-annealed condition contain carbon in

solution. When the steel is heated to the temperature range of 425 to 815 °C (800 to 1500 °F), as in the weld heat-affected zone (HAZ), carbon combines with chromium, resulting in chromium carbide precipitation at the grain boundaries. Precipitation of chromium carbide (sensitization) makes the material susceptible to intergranular corrosion. Sensitization to intergranular corrosion is influenced in austenitic stainless steel by (1) the carbon content, (2) the solubility of the carbon, (3) the presence of one or more stabilizing elements, (4) the proximity of the actual metal temperature to 1200 °F, and (5) the time period during which the metal is held in that temperature range. The first three factors are controllable only through selection of composition; the latter two are controllable by adjustment of welding conditions such as heat input, size of spot, production rate, and provisions for cooling. Martensitic stainless steels most often welded are types 403,410,414, and 431. These steels can be resistance welded

in the annealed, hardened or hardened-and-tempered condition. Regardless of prior condition, welding produces a hardened martensitic zone adjacent to the weld. The hardness of this zone depends mainly on carbon content, although it can be controlled to a degree by welding procedure. As the hardness of the metal in the HAZ increases, its susceptibility to cracking increases and its toughness decreases. Steels having a maximum carbon content of 0.15%, such as types 403 and 410, often produce satisfactory welds without postweld heat treatment. Steels with higher carbon content, such as types 420 and 440A, generally require postweld heat treatment. Ferritic grades of stainless steel most often welded are types 405,430,442, and 446. Resistance welds in these steels

exhibit reduced ductility at room temperature because of grain growth in the HAZ. Ferritic stainless steels can also become embrittled if heated to approximately 475 °C (885 °F). This phenomenon, known as "885 °F embrittlement," may be encountered in the HAZ of a ferritic stainless steel weld. If properties in the as-welded condition are unsuitable, annealing must follow welding. Postweld annealing, in addition to improving ductility, helps restore normal corrosion resistance. Precipitation-hardening grades of stainless steel can be resistance welded in the annealed or aged (hardened)

condition. If they are welded in the aged condition, the strength of the weld area is significantly reduced because the heat of welding anneals the material subjected to it. Generally, most of the strength may be restored by simply re-aging the weldment. When maximum strength is required in the weld area, it is best to resistance weld these grades in the annealed condition and then age the weldment. Procedure Development and Practice Considerations for Resistance Welding

Resistance Welding of Aluminum Alloys Aluminum alloys, both of the non-heat-treatable and heat-treatable types, either wrought or cast, can be resistance welded. Some of these alloys are welded more readily than others. Gas shielding may be required to maximize joint quality. Characteristics of aluminum alloys include comparatively high thermal and electrical conductivity, a relatively narrow plastic range (about 95 to 205 °C, or 200 to 400 °F, temperature differential between softening and melting), considerable shrinkage during cooling, a troublesome surface oxide, and an affinity for copper electrode materials. Base-Metal Characteristics

Although aluminum alloys can be resistance spot and seam welded, some alloys or combinations of alloys have higher aswelded properties than others. Table 15 gives melting ranges, electrical and thermal conductivities, and resistance weldability of some wrought alloys and casting alloys.

TABLE 15 MELTING RANGES, ELECTRICAL AND THERMAL CONDUCTIVITIES, AND RESISTANCE WELDABILITY OF COMMON ALUMINUM ALLOYS Wrought and casting alloys are identified by Aluminum Association designations.

ELECTRICAL ALLOY AND MELTING RELATIVE CONDUCTIVITY, THERMAL TEMPER RANGE CONDUCTIVITY(B),% %IACS(A) °C °F NON-HEAT-TREATABLE WROUGHT ALUMINUM ALLOYS 1350-H19 646119562 60 657 1215 1060-H18 646119561 57 657 1215 1100-H18 643119057 55 657 1215 3003-H18 643119040 39 654 1210 3004-H38 629116542 42 652 1205 5005-H38 632117052 51 652 1205 5050-H38 627116050 49 652 1205 5052-H38 593110035 35 649 1200 5083-H321 574106529 30 638 1180 5086-H34 584108431 32 640 1184 5154-H38 593110032 32 643 1190 5182-O 574106531 31 640 1185 5454-H34 602111534 34 646 1195 5456-H321 571106029 30 638 1180 HEAT-TREATABLE WROUGHT ALUMINUM ALLOYS 2014-T6 51095040 39 638 1180 2024-T361 50193530 31 638 1180 2036-T4 554103041 40 649 1200 2219-T37 543101028 29 643 1190 6009-T4 560104044 43 649 1200 6010-T4 585108539 38

RESISTANCE WELDABILITY(C)

ST ST RW RW RW RW RW RW RW RW RW RW RW RW

ST ST RW ST RW RW

649 1200 6061-T6 5931100649 1200 6063-T6 6151140654 1210 6101-T6 6151140652 1205 7075-T6 477890638 1180 ALUMINUM CASTING ALLOYS 413.0-F 5741065582 1080 443.0-F 5741065632 1170 308.0-F 521970613 1135 238.0-F 507945599 1110 513.0-F 5791075638 1180 520.0-T4 449840604 1120 333.0-T6 516960585 1085 C355.0-T61 5461015621 1150 356.0-T6 5571035613 1135 712-F 6041120643 1190

43

43

RW

53

51

RW

57

55

RW

33

33

ST

31

32

LW

37

37

RW

37

37

ST

25

26

LW

34

34

ST

21

22

NR

29

30

ST

39

38

ST

39

38

ST

40

39

RW

(A) INTERNATIONAL ANNEALED COPPER STANDARD. VOLUME BASIS AT 20 °C (68 °F). FOR COMPARISON, COPPER ALLOY 102 (OXYGEN-FREE COPPER) IS 101% AND LOW-CARBON (1010) STEEL ABOUT 14%. (B) BASED ON COPPER ALLOY 102 AS 100%. WHICH HAS A THERMAL CONDUCTIVITY OF 391 W/M · K (226 BTU/FT · H · °F) AT 20 °C (68 °F). LOW CARBON STEEL HAS A THERMAL CONDUCTIVITY OF ABOUT 13% ON THIS RELATIVE SCALE. (C) RW, READILY WELDABLE; ST, WELDABLE IN MOST APPLICATIONS BUT MAY REQUIRE SPECIAL TECHNIQUES FOR SPECIFIC APPLICATIONS; LW. LIMITED WELDABILITY AND USUALLY REQUIRES SPECIAL TECHNIQUES; NR. WELDING NOT RECOMMENDED Alclad Alloys. Resistance welding is also done on Alclad products made by roll cladding some of the alloys listed in

Table 15 with a thin layer of aluminum or aluminum alloy. Because this layer is anodic to the core alloy, it provides electrochemical protection for exposed areas of the core. Alclad alloys 2219, 3003, 3004, 6061, and 7075 have a cladding of alloy 7072, which contains 1% Zn; Alclad alloy 2014 has a cladding of alloy 6003 or sometimes alloy 6053, both of which contain about 1.2% Mg; and Alclad alloy 2024 has a cladding of alloy 1230, which contains a minimum of 99.3% Al. Effects on Weldability. The hardness of an alloy is one variable influencing weldability. Any alloy in the annealed

condition (O temper) is more difficult to weld than the same alloy in a harder temper. In general, alloys in the softer tempers are much more susceptible to excessive indentation and sheet separation and to low or inconsistent weld strength. Greater deformation under the welding force causes an increase in the contact area and variations in the distribution of current and pressure. Therefore, welding of aluminum alloys in the annealed condition or in the softer tempers is not recommended without special electromechanical or electronic controls.

High-strength alloys such as 2024 and 7075 are easy to resistance weld, but may require applications of a forge pressure, because they are more susceptible to cracking and porosity than the lower-strength alloys. Resistance Welding Machines Aluminum alloys can be resistance welded with single-phase direct-energy, three-phase direct-energy, and stored-energy machines. Best results are obtained by using a machine that has these features: • • • • •

ABILITY TO HANDLE HIGH WELDING CURRENTS FOR SHORT WELD TIMES SYNCHRONOUS ELECTRONIC CONTROLS FOR WELD TIME AND WELDING CURRENT LOW-INERTIA WELDING HEAD FOR RAPID FOLLOW-UP OF ELECTRODE FORCE SLOPE CONTROL (FOR SINGLE-PHASE WELDING MACHINES) MULTIPLE-ELECTRODE-FORCE SYSTEM TO PERMIT PROPER FORGING OF THE WELD NUGGET AND REDRESSING OF ELECTRODES

Electrodes and Electrode Holders Selection of electrode material and face shape, maintenance of the face, and cooling of the electrode are important in producing consistent spot and seam welds in aluminum alloys. Copper Alloy Electrodes. Resistance Welder Manufacturers Association (RWMA) classes 1, 2, and 3 are used for

welding aluminum alloys. These electrode materials have high electrical and thermal conductivities, which combined with adequate cooling help keep the temperature of the electrode below the temperature at which aluminum will alloy with copper and cause electrode pickup. Design of spot welding electrodes suitable for spot welding of aluminum includes both straight and offset electrodes.

Construction details of each are shown in Fig. 26. Straight electrodes should be used whenever possible, because deflection and skidding may occur with offset electrodes under similar welding conditions. If offset electrodes are used, the amount of offset should be the minimum permitted by the shape of the assembly being welded.

FIG. 26 CUTAWAY VIEW SHOWING INTERNAL COMPONENTS OF STRAIGHT AND OFFSET RADIUS-FACE ELECTRODES USED IN RESISTANCE SPOT WELDING OF ALUMINUM ALLOYS

Only electrodes that have the cooling-water hole within 9.5 mm (

3 in.) of the face surface should be used. 8

Preweld Surface Preparation Although welds for some purposes can be made satisfactorily without any preweld surface preparation, welds that are free of cracks, porosity and sheet separation, and that have the most uniform strength and symmetry, are obtained only with correct procedures for cleaning and reduction or removal of oxide film. In addition, adequate surface preparation reduces electrode contamination. Procedure Development and Practice Considerations for Resistance Welding

Resistance Welding of Copper and Copper Alloys Resistance spot welding is widely used for joining copper and copper alloys. Principal applications include welding sections up to about 1.52 mm (0.060 in.) thick, particularly those alloys with low electrical conductivities. Many copper alloys with low conductivities can be seam welded easily. Coppers with a low weldability factor are difficult to seam weld. Projection welding is not recommended for copper or for most brasses. Bronzes can be projection welded with satisfactory results in many applications. Welding Characteristics The resistance weldability of any copper alloy is inversely proportional to its electrical and thermal conductivities. Generally, alloys with lower conductivities are easier to weld. Compared with steel, most copper alloys require shorter weld time, lower electrode force, higher current, and different electrode materials that are compatible with the alloy being welded. Welding Equipment Single-phase and three-phase direct-energy and electrostatic stored-energy (capacitor-discharge) welding machines are used for resistance welding of copper and copper alloys. The addition of slope control to single-phase direct-energy welding machines is not necessary for spot welding most copper alloys. In welding high-zinc brasses, the use of upslope can result in an increase of as much as 20% in weld strength. Downslope is not recommended for welding of any of the copper alloys. Welding Machine Controls. Copper alloys are particularly sensitive to variations in welding conditions, and therefore all direct-energy machines used for welding these alloys should be equipped with synchronous electronic controls, especially in applications requiring short weld times. These devices are capable of controlling weld time and welding current for repeated operations with extreme accuracy.

Electrodes The current used for resistance welding of copper alloys is much higher than that used for welding low-carbon steel, and therefore, the electrode must have high electrical conductivity to minimize heat buildup. Electrode Materials. The Resistance Welder Manufacturers Association (RWMA) class 1 electrode materials (typically

tungsten or molybdenum alloys), containing copper and cadmium, are sometimes used for welding copper and highconductivity brass and bronze. Class 2 materials, containing copper and chromium, are used on low-conductivity brass and bronze and the copper-nickel alloys. Class 3 materials are used in electrodes for seam welding. Electrodes must be sufficiently water cooled to minimize sticking to the work metal and to prolong their life. Face contours must be carefully prepared and the electrodes must be properly aligned for welding. Selection of Process

Weldability of the work metal often determines which process should be used for a given application. Some of the coppers and copper alloys can be spot welded, but not seam welded because of high conductivity, and not projection welded because of low compressive strength of the projections at elevated temperature. Coppers Coppers and copper alloys having electrical conductivity higher than about 30% IACS are the least well suited for resistance spot, projection, or seam welding, mainly because of severe electrode pickup. Thin copper stock can be welded using electrodes faced with RWMA class 13 (tungsten) or class 14 (molybdenum), but surface appearance is poor and frequent electrode maintenance is required. A tinned coating on wire or sheet is helpful in welding copper. Beryllium Copper Beryllium copper alloys can be resistance welded most successfully in thin gages. Spot welding produces satisfactory welds; seam welding is less successful. Projection welding is satisfactory, provided that the projections can be formed with the work metal in the annealed condition and without cracking the work metal around the projection. Close control of welding conditions is required for consistent weld size and joint strength. Oxide films produced by heat treating must be removed to ensure low and consistent contact resistance. Work metals that have not been heated after rolling frequently need only degreasing before welding. Low electrical conductivity (22% IACS for alloys C17000 and C17200) contributes to the weldability of beryllium copper alloys. However, they are more difficult to resistance weld than low-carbon steel. Alloy C17500 has an electrical conductivity of 45% IACS and is more difficult to resistance weld than higher-strength, lower-conductivity beryllium copper. Low- and High-Zinc Brasses The low-zinc brasses are difficult to weld, although easier than copper, and are subject to electrode pickup. Welds made in these brasses may lack strength, principally because of comparatively high electrical conductivity (32 to 56% IACS). The high-zinc brasses have an electrical conductivity of 27 to 28% IACS and can be both spot and projection welded over a wide range of conditions. Electrode pickup can be a problem unless weld time, welding current, and electrode force are properly selected. Excessive electrode pickup and blowthrough of the weld may occur when long weld times, high energy input, and low electrode forces are used. Yellow brasses (alloys C26800 and C27000) are less susceptible to electrode pickup than cartridge brass except when long weld times and high energy input are used. Electrode force should be sufficient to prevent arcing or expulsion of molten metal, to which these alloys are subject because of their 30 to 40% Zn content, which boils at about 905 °C (1665 °F). As shown in Table 16, the recommended electrode force, when using electrodes having a face diameter of 4.8 mm (

3 in.), is approximately 1.8 kN (400 lbf). 16

TABLE 16 GUIDELINES FOR RESISTANCE SPOT WELDING OF SELECTED COPPER ALLOYS

UNS NO.

C23000 C24000 C26000 C26800-C27000 C28000

ALLOY NAME

RED BRASS LOW BRASS CARTRIDGE BRASS YELLOW BRASS MUNTZ METAL

WELDING PARAMETER(A) WELD TIME, ELECTRODE FORCE CYCLES KN LBF 6 1.8 400 6 1.8 400 4 1.8 400 4 1.8 400 4 1.8 400

WELDING CURRENT,A 25,000 24,000 25,000 24,000 21,000

C51000-C52400 C62800 C65100-C65500 C66700 C68700 C69200

PHOSPHOR BRONZE ALUMINUM BRONZE SILICON BRONZE MANGANESE BRASS ALUMINUM BRASS SILICON BRASS

6 4 6 6 4 6

2.3 2.3 1.8 1.8 1.8 2.3

510 510 400 400 400 510

19,500 21,000 16,500 22,000 24,000 22,000

(A) FOR SPOT WELDING 0.91 MM (0.036 IN.) THICK SHEET USING RWMA TYPE E ELECTRODES WITH 4.8 MM (

3 16

IN.) DIAM FACE AND 30 ° CHAMFER AND MADE OF RWMA CLASS 1

MATERIAL Copper Nickels The copper-nickel alloys have electrical conductivities of 4.5 to 11% IACS, are readily spot and seam welded with relatively low welding current, and generally do not alloy with the electrode material and cause electrode pickup. Bronzes The phosphor bronzes, except alloy C50500, which is not recommended for resistance spot and seam welding because of its high electrical conductivity (48% IACS), have relatively low electrical conductivity (11 to 20% IACS) and am readily spot and seam welded using low welding currents. Electrode pickup can be reduced by use of a type F (radius) electrode face and frequent redressing to keep the face clean and smooth. Hot shortness can be minimized by supporting the workpieces to prevent strain during welding and by using a greater minimum overlap than recommended by the data in Table 17.

TABLE 17 GUIDELINES FOR RESISTANCE SPOT WELDING OF HIGH-ZINC BRASSES

THICKNESS OF THINNEST SHEET mm in. 0.81 0.032

MINIMUM SPOT SPACING mm in. 5 16

1.27

0.050

16

1.63

0.064

19

2.39 3.18

0,094 0.125

25 38

8 5 8 3 4

16 19

1 1

MINIMUM CONTACTING OVERLAP(A) mm in. 1 13

1 2

25 32

2 5 8 3 4

1 1

1 4

SHEAR LOAD OF JOINT lbf kn 1.47 330 2.28

512

3.02

680

5.20 8.33

1168 1872

(A) MINIMUM EDGE DISTANCE IS EQUAL TO ONE HALF THE CONTACTING OVERLAP. Procedure Development and Practice Considerations for Resistance Welding

Nickel Silvers Nickel silvers, which have about the same conductivities (6 to 10.9% IACS) as copper nickels, are spot welded as readily as copper nickels but are more difficult to seam weld. Surface contaminants such as lead and bismuth (which form lowmelting eutectics with copper and nickel) or sulfur (which may be introduced in forming) must be removed before resistance welding. Procedure Development and Practice Considerations for Resistance Welding

References

1. L.P. CONNER, ED., WELDING HANDBOOK, 8TH ED., VOL 1, AWS, 1987, P 16 2. J.F. YOUNG AND R.S. SHANE, ED., MATERIALS AND PROCESSES, PART B: PROCESSES, 3RD ED., MARCEL DEKKER, 1985, P 11257-1126 3. J.G. BRALLA, ED., HANDBOOK OF PRODUCT DESIGN FOR MANUFACTURING, MCGRAW-HILL, 1986, P 7-35 4. J.F. YOUNG AND R.S. SHANE, ED., MATERIALS AND PROCESSES, PART B: PROCESSES, 3RD ED., MARCEL DEKKER, 1985, P 1127 5. J.G. BRALLA, ED., HANDBOOK OF PRODUCT DESIGN FOR MANUFACTURING, MCGRAW-HILL, 1986, P 7-38 6. R.L. O'BRIEN, ED., WELDING HANDBOOK, 8TH ED., VOL 2, AWS, 1991, P 552 7. J.G. BRALLA, ED., HANDBOOK OF PRODUCT DESIGN FOR MANUFACTURING, MCGRAW-HILL, 1986, P 7-39 8. "RECOMMENDED PRACTICES FOR RESISTANCE WELDING," AWS C1.1-66, AWS, 1966 9. RESISTANCE WELDING MANUAL, 3RD ED., VOL 1, RESISTANCE WELDER MANUFACTURERS ASSOCIATION, 1956, P 88 10. R.L. O'BRIEN, ED., WELDING HANDBOOK, 8TH ED., VOL 2, AWS, 1991, P 593 11. H.B. CARY, MODERN WELDING TECHNOLOGY, 2ND ED., PRENTICE-HALL, 1989, P 251-252 Procedure Development and Practice Considerations for Electron-Beam Welding

Introduction ELECTRON-BEAM WELDING (EBW) can produce deep, narrow, and almost parallel-sided welds with low total heat input and relatively narrow heat-affected zones in a wide variety of common and exotic metals. In many applications, a single pass is sufficient to weld the components. When initially applied to industrial applications in the 1950s, EBW had to be performed totally in a vacuum environment. In this form, EBW was primarily used because of its ability to reliably produce very high quality welds in materials normally considered extremely difficult to join (for example, refractory metals, oxidizable metals, stainless steels, and superalloys). This made it ideally suited for use in aerospace, nuclear, and other similar (high-technology) manufacturing industries. During the 1960s, the capability of employing the EBW process for medium-vacuum and nonvacuum welding tasks was developed. This increased production capacity helped make EBW more attractive to other industries desiring to take advantage of the numerous joining benefits it provided. However, the use of EBW required that all applications have a demonstrated capacity for fairly high-volume production. Today, EBW is used to join not only difficult-to-weld materials but also common materials such as steel ( ≤ 150 mm, or 6 in., thick), aluminum ( ≤ 305 mm, or 12 in., thick), and copper ( ≤ 100 mm, or 4 in., thick) (Ref 1). These and other metals (in thicknesses ranging from foil to plate dimensions) can be electron-beam welded. In addition, EBW can be used to join numerous combinations of dissimilar metals. See the Section "Fusion Welding Processes" in this Handbook for a description of EBW systems and equipment.

Reference

1. A. SANDERSON AND K.R. NIGHTINGALE, WELD. J., APRIL 1990, P 45 Procedure Development and Practice Considerations for Electron-Beam Welding

Weld Geometry Part Configuration. Whether a component is welded in a vacuum or at atmospheric pressure (nonvacuum), part

configuration and the corresponding joint designs used are critical parameters in the successful application of EBW. A circular weld in the axial direction, for example, that joins a disk-shaped member (gear or cup) to a shaft (Fig. 1) may

experience severe constraint. Such joints are practical in easily welded metals--those that are soft, have low yield points, and exhibit low shrinkage--but not in difficult-to-weld steels that have high strength, hardenability, or machinability due to high sulfur content.

FIG. 1 CUTAWAY VIEW SHOWING LOCATION AND PENETRATION OF A CIRCULAR ELECTRON-BEAM WELD JOINING A DISK-SHAPED COMPONENT TO A SHAFT

The difficulties encountered in welding shapes such as those shown in Fig. 1 are aggravated if the weld zone is broadened by defocusing the electron beam. This is a frequent problem if exact beam position cannot be controlled, either because of beam instability, magnetism of the parts, or excessive tolerances in fixturing on high throughput machines. The need for beam broadening can be eliminated by the use of closed-loop beam position control systems. Shrinkage stress is minimal in radial welds, as shown in Fig. 2. The volume of molten metal is less constrained by the axial direction if the gear hub is allowed to move by using push fit and no restraints. The weld must also have root clearances--the weld zone must not extend into the shaft. Stresses are lower if the melt zone has parallel side boundaries. Triangular weld zones that are caused by broadened beams, uneven beam power distribution, or incomplete penetration have more complex stresses which may result in cracking.

FIG. 2 SHRINKAGE STRESS IN A RADIAL WELD

The same principles of joint design discussed above for joints between shafts and disklike members apply equally to Tjoints and comer joints between plates. Figure 3 shows recommended and nonrecommended weld configurations. Although true for all welding methods, it is essential in EBW (which is usually performed in a single pass) that constraints on the volume of the melt must be avoided by self-aligning interlocking steps in the joint. In multilayer arc welds, the individual weld bead is less constrained in volume and shrinks from its surface downward. Functionally, it is not useful to make the cross section of the joint any less thick than the plates that are being joined. If two welds are made in succession, the second weld may be constrained by the first (Fig. 3), which may lead to solidification cracking. It is thus advisable to direct the beam at the joint parallel to the faying surfaces and to cause melting over the whole contact area.

FIG. 3 OPTIMUM VERSUS LEAST DESIRABLE WELD CONFIGURATIONS. (A) NOT RECOMMENDED--MAXIMUM CONFINEMENT OF MOLTEN METAL, MINIMUM JOINING CROSS SECTION (ARROWS); WASTES BEAM ENERGY FOR MELTING, NONFUNCTIONAL METAL. (B) MOST FAVORABLE--VOLUME OF MELT NOT CONFINED; MAXIMUM JOINING CROSS SECTION (ARROWS). (C) NOT RECOMMENDED--MAXIMUM CONFINEMENT OF MELT (UNLESS GAP IS PROVIDED); JOINING CROSS SECTION LESS THAN PLATE CROSS SECTION. (D) MOST FAVORABLE-MINIMUM CONSTRAINT AND CONFINEMENT OF MELT; MINIMUM INTERNAL STRESSES; WARPAGE CAN BE OFFSET BY BENDING PRIOR TO WELDING; TILT CAN BE OFFSET BY LOCATION OF T-ARM AT LESS THAN 90°

TO BASE PRIOR TO WELDING. FILLET OBTAINED BY PLACING WIRE IN RIGHT CORNER AND MELTING IT WITH THE BEAM. (E) NOT RECOMMENDED--TWO SUCCESSIVE WELDS; SECOND WELD IS FULLY CONSTRAINED BY THE FIRST WELD AND SHOWS STRONG TENDENCY TO CRACK.

Surface Geometry. Usually, EBW does not use or need filler wire. Therefore, neither V-grooves nor corners can be

filled; in fact, they are detrimental. Surface tension pulls some of the molten metal up the V-shaped surfaces so that it is no longer available to fill the space created by the keyhole that is generated by the moving beam. Voids remain in the actual joint plane (Fig. 4).

FIG. 4 GAP IN JOINT PLANE CREATED BY SURFACE TENSION

A step in the surfaces at the joint line is also undesirable. Any small lateral shift of the beam from the low to the high side, and vice versa, changes the penetration to some extent. In the extreme case, this shift may be equal to the height of the step, as shown in Fig. 5. On circular welds, with joint configurations similar to Fig. 5(a), a noncircular weld profile may develop with irregularities that can lead to uneven applied stresses. By contrast, a mere shift in the melt zone, as shown in Fig. 5(b), does not affect weld quality, as long as the entire joint line is covered.

FIG. 5 EFFECT OF BEAM AND MELT ZONE POSITION ON WELD PENETRATION. (A) INCLINED ELECTRON BEAM. (B) VERTICAL ELECTRON BEAM

Configurations for Wide Welds Bridging a Gap. The high speed and autogenous nature of an electron-beam weld can 1 1 be advantageous when a wide weld (6.5 to 13 mm, or to in.) is needed without great depth-to-width ratio. The 4 2

atmospheric electron beam has been used extensively to make such welds in mass-production automotive parts such as catalytic converter cases and die-cast aluminum intake manifolds. The easiest weld to make with an electron beam is an edge-weld in the vertical position of the beam (weld horizontal and flat), as illustrated by the edge-flange weld in Fig. 6. In vacuum, the beam must be defocused; in air, the standoff must be short enough to prevent substantial beam scattering. The standoff is not critical because energy flow is not influenced by standoff as it is in electric arc welding.

FIG. 6 EDGE-FLANGE WELD TYPICALLY USED TO JOIN THIN PANELS

For edge welds, changes in energy input per unit length and changes in power density affect only the amount of molten metal. Offset and gaps up to the thickness of the sheet to be joined often can be tolerated (Ref 2). In making an edgeflange weld (Fig. 6), the same advantages exist, provided the flange is high enough. For joining thinner sheet that is tightly clamped, edge welds also can be made with the electron beam in the horizontal position. The nonvacuum EBW (EBW-NV) process can tolerate and fill large gaps if sufficient sacrificial metal is available to be melted by the beam (Fig. 7). If not, surface tension may pull whatever metal is melted to either side of the weld joint.

FIG. 7 USE OF SACRIFICIAL METAL TO FILL A LARGE GAP IN A SQUARE-GROOVE WELD

Melt-Zone Configuration. Welds with parallel sides of the melt zone (Fig. 8a) show minimum warpage and only

transverse shrinkage when they are unconstrained. Triangular melt zones (Fig. 8b) cause warpage. If there are constraints, the hindered warpage and shrinkage cause internal stresses. A parallel-sided melt zone is preferred in either case; to obtain this configuration, joint design must allow for a full-penetration weld.

FIG. 8 VARIATION IN BEAM VELOCITY AND SUBSEQUENT WARPAGE OF 0.25 MM (0.10 IN.) THICK TYPE 304 STAINLESS STEEL JOINED BY A BUTT WELD. (A) BEAM TRAVERSE, 7110 MM/MIN (280 IN./MIN); TRIANGULAR MELT ZONE. (B) BEAM TRAVERSE, 4190 MM/MIN (165 IN./MIN); PARALLEL-SIDED MELT ZONE. WELD PARAMETERS: 12 KW BEAM POWER; 11.4 MM (0.45 IN.) STANDOFF; HELIUM SHIELDING GAS


2. J.F. LOWRY, J.H. FINK, AND B.W. SCHUMACHER, A MAJOR ADVANCE IN HIGH-POWER ELECTRON BEAM WELDING IN AIR, J. APPL. PHYS., VOL 47, 1976, P 95-106 Procedure Development and Practice Considerations for Electron-Beam Welding

Joint Design Figures 9, 10, 11, 12, and 13 show the commonly used joint and weld types for EBW. Different joint preparations, joint designs, and welding positions are used to meet special requirements. For preferred joint configurations and shrinkage stresses encountered in various joint designs, see the section on weld geometry for electron-beam welding in this article.

FIG. 9 TYPES OF BUTT JOINTS AND WELDS TYPICALLY GENERATED BY EBW

FIG. 10 TYPES OF CORNER JOINTS AND WELDS TYPICALLY OBTAINED WITH EBW

FIG. 11 TYPICAL WELDS PRODUCED BY EBW OF T-JOINTS

FIG. 12 TYPICAL WELDS OBTAINED BY EBW OF LAP JOINTS

FIG. 13 TYPICAL WELDS GENERATED BY EBW OF EDGE JOINTS HAVING COMPONENTS OF EQUAL AND UNEQUAL SECTION THICKNESSES

Welds in Butt Joints Butt joints are the most common of the five basic joint types used in EBW. The basic square-groove weld, which is shown in Fig. 9(a), 9(b), and 9(c), needs only the simplest and least expensive joint preparation. It is suitable for either partial- or full-penetration welds. Good fit-up and external fixtures are needed. The flush joint (Fig. 9b) is preferred to the stepped joint (Fig. 9c) for joining unequal thicknesses, chiefly because control of conditions for making sound fullpenetration welds is less critical than for the stepped joint. A wider beam is used in welding the stepped joint, and the beam angle must he carefully controlled to avoid scarfing the upper edge of the thicker member (angle too small) or missing the bottom of the joint (angle too large), especially if the thinner member is more than about 9.5 mm (

3 in.) 8

thick. In Fig. 9(d), the joint is rabbeted to make it self-aligning, but the offset is small to avoid leaving an unwelded seam near the root of the weld (compare Fig. 9e). In Fig. 9(d) to (g), the joints are self-aligning and may be self-fixturing in circular, circumferential, and certain other joint arrangements. In Fig. 9(e) and 9(f), the joints are both self-aligning and selfbacking; each, however, leaves an unwelded seam near the root of the weld. In Fig. 9(g) and 9(h), two ways are shown of providing integral and filler metal. The lip of the joint in Fig. 9(g) provides more filler metal than the shoulder of the joint in Fig. 9(h), but because it conceals the joint, a scribed line or other means for beam placement and for scanning must be provided before the welding operation can proceed. The slant-groove weld (also called angular or scarf) shown in Fig. 9(j) is used in butt joints to facilitate fixturing. It is also used where limitations on the beam location prevent fusion for the entire groove depth if the basic square-groove weld (Fig. 9a) were used. Greater weld depth, as measured parallel to the beam axis, is needed for full penetration of a given metal thickness than with the square-groove weld. This type of weld and variations of it are used both in butt joints and in corner joints. A slant-butt joint can also be welded with beam alignment set at 90° to the surface of the work. This beam angle permits the production of defect-free welds where fit-up is poor or where the joint opening is larger than can be tolerated when the beam is aligned with the groove. In a related type of weld, the joint has a square-groove preparation, and the beam is at an angle to the groove. The main reason for using this type of weld is difficulty of access with other types of welds. Welds in Corner Joints Eight types of electron-beam welds frequently made in corner joints are shown in Fig. 10. Corner joints are second only to butt joints in frequency of use for EBW.

Important differences between butt and corner joints are notch sensitivity and suitability for nondestructive testing. To compare alignment, self-fixturing, and the occurrence of portions of unwelded seam, see Fig. 9 and 10. The two simplest and most economical welds in corner joints are the melt-through (Fig. 10a) and basic square-groove (Fig. 10b) welds. Neither is self-aligning or self-fixturing; manipulation of the work for weld (b) in Fig. 10 can be simplified by using a horizontal electron beam and corresponding work orientation (rotated 90° from the orientation shown). The melt-through weld (Fig. 10a), unless made with a fusion zone wide enough to eliminate the unwelded seam completely, is weaker than the square-groove weld and is notch sensitive. The corner-flange weld (Fig. 10h), usually made only on thin stock, requires precision forming of the 90° bend. Welds in corner joints are subject to high stress concentrations. Welds in T-Joints Three types of electron-beam welds made in T-joints are shown in Fig. 11. The melt-through or blind weld (Fig. 11a), like melt-through welds in butt or corner joints, leaves an unwelded seam with resulting low strength, notch sensitivity, and corrosion susceptibility. Figure 11(b) is the preferred joint design. The double-square-groove weld (Fig. 11c) is used primarily on sections 25 mm (1 in.) or more thick. For welds similar to Fig. 11(b) and 11(c), filler metal can be added. Welds in Lap Joints Three types of electron-beam welds made in edge joints are shown in Fig. 13. Thick sections can be joined by deep, narrow square-groove welds (Fig. 13a). Shallow welds made with a low-power, partially defocused beam are used to join thin sections to each other or to thick sections, as shown in Fig. 13(b) and 13(c). Shallow edge welds can be made at high speed. Welds of this type are particularly useful on hermetically sealed assemblies, which may be designed for planned salvage by removal of the weld region and subsequent rewelding. Butt Joints Versus Corner and T-Joints The concentration of bending stresses at the weld in corner and T-joints is sometimes avoided by designing the work for a butt joint on a straight section near the corner or T (flange neck or welding neck construction). This principle of design is applicable also to butt joints between members having different thicknesses, and it can be applied to straight line, circular, and circumferential joints. In welding such a joint in magnetic metals, if the beam path is too close to the corner, T, or shoulder, the beam may be deflected away from the metal that it passes near on its way to the joint, causing it to enter the work at an angle and possibly to miss the lower portion of the joint. To avoid this, the beam clearance should be 0.8 mm (0.030 in.) or more, depending on the height of the magnetic projection (Fig. 14).

FIG. 14 JOINT LOCATION AND BEAM CLEARANCE FOR AN EBW BUTT JOINT NEAR A PROJECTING CORNER, T, OR SHOULDER CONSISTING OF MAGNETIC WORK METAL

Procedure Development and Practice Considerations for Electron-Beam Welding

Special Joints and Welds Many variations of the basic joint and weld types described in the preceding section and illustrated in Fig. 9, 10, 11, 12, and 13 are used to meet the needs of individual applications. Some of these special joints and welds are discussed below. Plug and puddle welds are usually made by manually (or automatically) manipulating the work under a fixed beam at

low power. Filler metal used for plug welds is often preplaced at the weld site. Puddle welding is used mainly to fuse shallow defects together locally. Multiple-Pass Welds. Most electron-beam welds produce the desired penetration in a single pass. Tack welds and

cosmetic or smoothing passes are not considered penetration passes. Welds several inches deep can be made in most metals in a single pass. Weld depth obtainable in a single pass can be nearly doubled by welding from both the face and the back of the work. In a variation on straight-through two-pass welding from opposite sides, separate welds that meet or almost meet can be made at an angle of 90° to each other in a rabbeted square-groove joint. Tangent-tube welds are longitudinal welds joining two parallel tangent tubes (or cylinders). The tubes may differ in

size; the beam is perpendicular, or nearly perpendicular, to the common axial plane of the two tubes. Added filler metal may be used for reinforcement. Thinner wall tubes can be joined more easily by EBW than by arc welding methods. Three-Piece Welds. Welds can be made that join three or more pieces in which there is penetration in a single pass into

all of the pieces. Many of the difficulties encountered in welding very thin metal or foil are eliminated by sandwiching it between two thicker sections. Multiple-tier welds are welds made simultaneously in in-line, separated joints (usually butt joints) in a single pass of the

electron beam. A more detailed discussion is given in the section of this article on multiple-tier welding . Welds using integral filler metal may be of the types shown in Fig. 9(g) and 9(h), in which an overhanging lip or a

shoulder provides filler metal to a butt joint, or both members may be made thicker at the joint than elsewhere.


EBW in High Vacuum Most EBW is done in a vacuum environment where the maximum ambient pressure is less than 0.13 Pa (1 × 10-3 torr). Maintenance of this degree of vacuum is important because of the effect that ambient pressure has on both the beam and the weld produced, as discussed below. Applications. High-precision applications that require welding to be done in a high-purity environment to avoid contamination by oxygen or nitrogen are ideally suited for high-vacuum EBW (EBW-HV). These applications include the nuclear, aircraft, missile, and electronics industries. Typical products include nuclear fuel elements, pressure vessels for rocket propulsion systems, special alloy jet engine components, and hermetically sealed vacuum devices. Welding of reactive metals such as zirconium and titanium in high vacuum is preferred over medium- or nonvacuum EBW because of the affinity of reactive metals to oxygen and nitrogen. Effect of Pressure on Beam. Under ambient high-vacuum conditions, the frequency of collisions between beam

electrons and residual gas molecules is extremely low, and any dispersion (i.e., beam-broadening effect) that would result from such scattering collisions is minimal. However, because the frequency of scattering collisions increases with the density of gas molecules present, this effect becomes greater as the surrounding environment pressure is increased. Only at ambient pressure values below 0.13 Pa (1 × 10-3 torr) is the effect of scattering insignificant enough that the beam can be held in sharp focus over distances of several feet (depending on the particular characteristics of the electron gun and electron optics being employed) to achieve maximum effectiveness in producing relatively deep, narrow welds. Beam scattering from collisions with residual gas molecules begins to occur at ambient pressures above 0.13 Pa (1 × 10-3 torr) and initially results in decreased penetration and increased weld width. This scattering effect reduces the maximum working distance that can be employed for producing a similar electron-beam weld as obtained under high-vacuum conditions at ambient pressures greater than 0.13 Pa (1 × 10-3 torr). In EBW-NV, where the ambient pressure is 10 MPa (760 torr) and above, the scattering dispersion of the beam increases so that penetration and working distance capacity are appreciably reduced below values obtainable in either high- or medium-vacuum processes. However, the penetration and working distance capabilities of EBW-NV are appreciably greater than those of conventional arc welding processes. Electron-beam dispersion characteristics and their effect on weld penetration at various ambient pressures are shown in Fig. 15.

FIG. 15 EFFECT OF PRESSURE ON BEAM INTENSITY AND RESULTING WELD QUALITY. (A) DISPERSION PATTERNS OF ELECTRON BEAM AT SELECTED PRESSURES. (B) WELD PENETRATION AS A FUNCTION OF AMBIENT PRESSURE LEVEL. ADDITIONAL PARAMETERS SUCH AS BEAM VOLTAGE, BEAM TRAVEL DISTANCE, AND TYPE OF AMBIENT GAS USED CAN PRODUCE A WIDE RANGE OF VALUES, THUS ACCOUNTING FOR THE SPREAD IN THE PENETRATION DATA.

Effect of Pressure on Weld and Heat-Affected Zone. The maintenance of a high vacuum in the work chamber

protects the weld metal and the HAZ from oxidation and contamination by harmful gases, serving the same function as inert shielding gases do in arc welding, while also degassing the weld. The contamination level as a function of pressure is discussed in the section EBW in Medium Vacuum in this article. The nearly complete absence of gaseous impurities eliminates the serious contamination difficulties that usually are encountered when attempting to arc weld relative metals such as titanium and zirconium. Width of Weld and Heat-Affected Zone. Electron-beam welds made in a high vacuum are narrower and have a

narrower HAZ than comparable welds made in medium vacuum or at atmospheric pressure, and they are much narrower than the narrowest welds made in production welding by GTAW. The narrow width of the heat-affected zone in

hardenable steels and other hardenable alloys permits welding of these metals after heat treatment, in many instances without loss of strength. Limitations. A major limitation in the use of high vacuum in the work chamber is the effect on unit production time,

because of the need to pump the chamber down before each load is welded. Pumpdown time is typically about 3 min for a 0.85 m3 (30 ft3) chamber and 10 min for an 8.5 m3 (300 ft3) chamber. The above pumpdown times are realistic only for a very clean and well-maintained system. In production, pumpdown may be nearly double the above times. The effect of this limitation on unit production time is reduced by welding a number of assemblies in each load and by keeping chamber size as small as possible. Small chambers specially designed for use with small workpieces have typical pumpdown times of I min (3.69 × 106 mm3, or 225 in.3, chamber) and 15 s or less (1.0 × 106 mm3, or 60 in.3, chamber). A second major limitation in the use of high vacuum is that work size is limited by chamber dimensions. This limitation is sometimes circumvented by the use of a chamber having special openings and seals that permit oversize work to extend outside the chamber, or by the use of a portable clamp-on chamber. Welding Conditions. Welding in high vacuum is done with low-voltage, as well as high-voltage, equipment. Beam

voltage ranges from 15 to 185 kV; beam current, 2 to 1000 mA; and welding speed, 100 to 5000 mm/min (4 to 200 in./min) for most applications. However, higher welding speeds are used for special thin section applications. The energy input to the work needed to produce a weld of the required depth of penetration and width is the basis for selecting the welding conditions. Depth of penetration is increased by increasing the voltage or current for greater beam power, or by decreasing the welding speed. Table 1 lists approximate energy inputs per inch of weld length for making narrow single-pass electronbeam welds 6.4 to 75 mm (0.25 to 3 in.) deep in the weldable alloys of copper, iron, nickel, aluminum, and magnesium. These values are intended to serve as guidelines for establishing conditions for welding work for which no previous experience is available. Energy-input requirements for specific applications depend on the alloy composition and special operating conditions such as the use of beam oscillation or a defocused beam.

TABLE 1 ENERGY INPUT AT THE WELD FOR SINGLE-PASS EBW IN HIGH VACUUM FOR VARIOUS DEPTHS OF PENETRATION

DEPTH OF PENETRATION mm in. 6.35 0.25 12.7 0.50 19.1 0.75 25.4 1.00 38.1 1.50 50.8 2.00 63.5 2.50 76.2 3.00

ENERGY INPUT PER WELD LENGTH FOR INDICATED METAL(A) Mg Cu Fe Ni Al kj/mm kj/in. kj/mm kj/in. kj/mm kj/in. kj/m kj/in. kj/mm 0.3 0.7 0.2 5 0.2 4 0.1 2 0.04 0.6 0.15 0.4 10 0.3 8 0.2 5 0.1 1.0 0.25 0.7 18 0.6 15 0.3 8 0.2 1.5 0.37 1.1 27 0.9 22 0.5 13 0.3 2.4 0.62 1.8 46 1.5 39 0.9 23 0.6 3.4 0.87 2.7 68 2.4 60 1.4 35 0.9 4.4 112 3.5 90 3.1 80 1.9 47 1.1 5.4 137 4.4 112 3.9 100 2.3 59 1.4

kj/in. 1 3 5 8 15 22 29 36

(A) VALUES ARE APPROXIMATE AND APPLY TO THE COMMONLY WELDED ALLOYS OF THE METALS. THEY ARE INTENDED TO SERVE AS STARTING POINTS FOR ESTABLISHING CONDITIONS FOR MAKING NARROW WELDS. ENERGY INPUT MAY VARY SUBSTANTIALLY FROM THESE VALUES IN SPECIFIC APPLICATIONS, DEPENDING ON THE COMPOSITION OF THE ALLOY AND SPECIAL OPERATING CONDITIONS. ENERGY INPUT AT WELD = [(BEAM VOLTAGE) (BEAM CURRENT) 0.6]/WELDING SPEED, WHERE ENERGY INPUT IS IN KJ/IN.; BEAM VOLTAGE IS IN KV; BEAM CURRENT IS IN MA; AND WELDING SPEED IS IN IN./MIN.


EBW in Medium Vacuum Electron-beam welding in a medium vacuum (EBW-MV) usually is performed at a welding (work) chamber pressure of about 10 Pa (7.5 × 10-2 torr), or at a work chamber pressure ranging from 0.4 to 40 Pa (3 × 10-3 to 3 × 10-1 torr)--the total partial-vacuum region. The gun chamber, where the beam is generated, is held at 1.3 to 13 mPa (10-5 to 10-4 torr), as in high-vacuum welding. A diffusion pump is required to provide this gun chamber vacuum, but only a mechanical-type pump is required to evacuate the work chamber to the final partial vacuum welding pressure desired. Comparison with High-Vacuum EBW. The chief advantage of welding in medium vacuum, compared with welding in

high vacuum, is the short pumpdown time for the welding chamber, which ordinarily does not exceed 40 s for a generalpurpose chamber (of about 0.11 m3, or 4 ft3) and which is less than 5 s for the specially designed chambers used in welding small parts. Accordingly, EBW-MV permits mass production of parts, using a chamber of minimum volume. Production rates for welding in medium vacuum depend on part design and other factors; maximum production rate is typically about 60 pieces per hour for general-purpose, manually operated equipment; upwards of 600 pieces per hour for specially tooled semiautomatic machines; and 1500 pieces per hour for fully automatic single-purpose machines. Generalpurpose chambers and tooling are used for short runs. Welds on reactive metals and other welds for which the effects of contamination by gases are especially critical are preferably made in high vacuum. The contamination level of air (total concentration of gases present) is proportional to pressure, as follows:

PRESSURE PA TORR 0.0013 10-5 0.13 10-3 13 10-1 50 4 × 10-1

GASES, PPM 0.01 1.3 132 500

Lower contamination levels are obtainable than those actually observed for arc welding with inert-gas shielding. Penetration and Weld Shape. Penetration at a given power level is ordinarily 5 to 10% less than for EBW-HV, and

welds are wider and slightly more tapered. These changes are caused by the scattering of the electrons in the beam by collision with gas molecules, which makes the beam broader. The effect of welding chamber pressure on penetration and weld shape is shown in Fig. 16 for welds made on type 304 stainless steel, using a beam length of 405 mm (16 in.) without changing focus. Maximum penetration shown in Fig. 16 is approximately 13 mm (

1 in.). Welding conditions were 150 kV, 30 mA, and 1525 mm/min (60 in./min). Penetration 2

drops off rapidly at pressures of 13 Pa (10-1 torr) or more at this beam length and at somewhat lower pressures when the beam path is longer.

FIG. 16 EFFECT OF WELDING-CHAMBER PRESSURE ON PENETRATION AND WELD SHAPE FOR TYPE 304 AUSTENITIC STAINLESS STEEL

Comparison With Nonvacuum EBW Welding. Disadvantages of medium-vacuum welding, compared with welding at atmospheric pressure, include limitations on work size imposed by the need to enclose the work in a vacuum chamber and the time needed for chamber pumpdown.

Beam deflection and oscillation, which cannot be used in welding at atmospheric pressure, can be used in EBW-MV. Applications. Most parts that can be welded in high vacuum can also be welded in medium vacuum. Depth of penetration is less than in high-vacuum welding, and depth-to-width ratio is less than about 10 to 1. Reactive metals are less readily welded in medium vacuum than in high vacuum. Medium-vacuum welding can be acceptable, however, for refractory metal welding when absorption of small amounts of oxygen can be tolerated. Medium-vacuum welding has been used for welding mass-produced automobile parts, such as gears and shafts, that are welded to close tolerances and that are not subsequently finished. Procedure Development and Practice Considerations for Electron-Beam Welding

Nonvacuum (Workpiece Out-of-Vacuum) Welding In EBW-NV (workpiece out-of-vacuum), which is also referred to as atmospheric EBW, the work to be welded is not enclosed in a vacuum chamber. Instead, it is welded at atmospheric pressure. A radiation-tight enclosure or chamber, similar to that used in high- or medium-vacuum EBW, surrounds the weld area. A vacuum, however, is not provided in this welding enclosure. A nonvacuum electron beam is generated at high vacuum in the same manner as in high- and medium-vacuum welding. It is focused down through a series of individually pumped stages, each connected to the other by concentrically aligned orifices, decreasing in diameter in the direction of increasing pressure; this differential pumping scheme provides the high vacuum (gun region) to atmospheric (workpiece region) pressure gradient necessary to allow the beam to pass down through the column and exit out into the ambient atmosphere. Figure 17 depicts a typical EBW-NV gun/column assembly. This unit can be mounted in either a vertical or a horizontal position and is capable of motion during welding. The gun/column assembly depicted in Fig. 17 is capable of beam power levels greater than 35 kW and beam transmission efficiencies in excess of 90%.

FIG. 17 SCHEMATIC WITH CUTAWAY VIEW SHOWING PRIMARY COMPONENTS OF A TYPICAL EBW-NV GUN/COLUMN APPARATUS. COURTESY OF LEYBOLD-HERAEUS VACUUM SYSTEMS INC.

Production rates are generally much higher for EBW-NV than for EBW-HV and EBW-MV, because no work chamber vacuum pumpdown time is required for each chamber load. This results in a lower per part production cost. Because pumpdown time for high- and medium-vacuum welding varies with chamber cleanliness and ambient humidity, elimination of work chamber pumpdown time in EBW-NV produces appreciable time savings. This saving, however, is gained at the expense of a reduction in penetration and working distance capability, thus imposing limits on the thickness and shape of the workpiece. Additionally, most (but not all) out-of-vacuum welds cannot be made as narrow and parallel sided as welds made in a vacuum. When welding in ambient atmospheric air alone, the practical working distance of the nonvacuum process is normally limited to the range of from 9.5 to 22 mm (

3 7 to in.); however, when welding in the presence of a lighter ambient gas, 8 8

such as helium, this practical working distance range is extended to 25 to 50 mm (1 to 2 in.). For this reason, most commercial nonvacuum EBW units are equipped to provide a "helium blowdown," which is an effluent of helium that

extends out from the nonvacuum column and along the beam travel path. The effect of this helium effluent is that it reduces the degree of beam scattering incurred using only air, thus allowing the working distance capability of the nonvacuum unit to be approximately doubled. Most metals that can be arc welded or electron-beam welded in a vacuum also can be welded in nonvacuum. As stated previously, inert-gas shielding of the weld zone can be provided when the molten weld metal (or HAZ) cannot be exposed to air. On readily welded metals, weld properties approach those of welds made in a vacuum; on difficult-to-weld metals, weld properties are usually lower than those of welds made in a vacuum. Ordinarily, filler metal is not used in EBW-NV. The addition of a suitable filler metal may be helpful, however, in joining metallurgically incompatible metals or in avoiding weld cracking caused by hot shortness. Selection of the filler-metal composition in such applications is critical. Operating conditions for nonvacuum welding differ substantially from those for welding in high or medium vacuum. Because beam dispersion increases in proportion to pressure and distance traveled, at pressures above the high-vacuum level, the nonvacuum beam becomes dispersed at travel distances greater than several inches in air at atmospheric pressure.

To provide an electron beam of sufficient energy to minimize the scattering effect of collisions with residual gas molecules within the electron-beam transfer column (i.e., the differentially pumped gun/column assembly), the applied gun voltage is held at a constant high value (from 150 to 200 kV). This also helps extend the work distance of the beam after it has exited the column and entered the ambient atmosphere. The practical working distance (standoff distance) for nonvacuum welding, as measured from the bottom of the final exit orifice to the top of the workpiece, ranges from 13 to 50 mm (

1 to 2 in.), although some nonvacuum welding is performed at both shorter and longer standoff distances. 2

With voltage held constant, beam current, working distance, and welding speed are selected to provide the required penetration and weld shape for a sound weld. Because it is usually desirable to keep working distance as short as possible and welding speed high, beam current is the primary control variable in nonvacuum welding. The small size of the exit orifice on the electron gun makes it necessary to focus the beam at or very close to the exit orifice. Accordingly, it is not possible to vary weld characteristics to a significant degree by changing focus. Similarly, beam deflection and oscillation are not possible. The absence of these adjustment capabilities in nonvacuum welding is not a handicap, because beam width is ordinarily great enough so that it does not promote undercutting or porosity at the root of the electron-beam weld. A cosmetic pass is seldom needed on electron-beam welds made at atmospheric pressure, because a smooth surface is produced on welds in most metals. The smoothness of the crown also improves the ease of nondestructive inspection of the as-welded surface; wire brushing is often adequate preparation for radiographic or liquid-penetrant inspection. The limitation on workpiece shape imposed by the need for a short working distance can be overcome to some degree by the use of electron gun nozzles specially designed to extend into restricted spaces. Filler metal is not ordinarily used except where necessary for weld reinforcement to produce the desired weld properties or to avoid cracking. A stream of dry filtered air, or an inert shielding gas such as argon or helium, is passed across the weld region in the space between the work and the electron gun, or may be supplied through a special insert nozzle assembly that is part of the exit orifice and is designed to minimize entrance of welding vapors and other contaminants into the gun. Auxiliary inert-gas shielding is supplied where needed for complete protection of the molten weld metal and the HAZ from gaseous contaminants. Welds are sometimes made on carbon and alloy steel and other readily weldable metals without using shielding gas. However, shielding gas is often used; helium is the preferred gas. Weld shape can be changed by varying the flow of helium. Penetration. Nonvacuum electron-beam welds are seldom deeper than 9.5 mm (

3 in.). Welds having penetration much 8

greater than 25 mm (1 in.) are easily produced, but ordinarily this requires a reduction in welding speed, with a corresponding increase in production cost. Figure 18 provides a comparison of penetration, as a function of weld speed, in steel and aluminum, employing a fixed power level and work distance and using a helium shield to improve power density.

FIG. 18 PLOT OF PENETRATION VERSUS WELDING SPEED AS A FUNCTION OF MATERIAL AT A BEAM POWER OF 9 KW

Studies made of nonvacuum welding of steel pipe showed that, at speeds of 2.5 to 25 m/min (100 to 1000 in./min), fullpenetration welds could be made in metal twice as thick if helium shielding gas was used instead of air. Table 2 gives the results of this comparison in terms of welding speed for joining steel pipe 1.3 to 3.8 mm (0.050 to 0.150 in.) thick in air or in helium. Welding speeds were 2.16 to 2.24 times greater in helium as in air. Beam power used was 12 kW.

TABLE 2 SPEEDS FOR NONVACUUM FULL-PENETRATION WELDING OF CARBON AND LOW-ALLOY STEEL PIPE AT A BEAM POWER OF 12 KW

WALL THICKNESS WELDING SPEEDS IN AIR IN HELIUM mm in. m/min in./min m/min in./min 1.27 0.050 10.6 417 22.9 900 2.03 0.80 6.12 241 13.4 528 2.54 0.100 4.75 187 10.5 412 3.81 0.150 3.00 118 6.68 264

RELATIVE WELDING SPEED IN HELIUM(A) 2.16 2.20 2.21 2.24

(A) BASED ON WELDING SPEED IN AIR BEING EQUIVALENT TO 1.00 Weld Shape and Heat Input. Nonvacuum electron-beam welds are generally wider and more tapered than highvacuum welds (medium-vacuum welds differ only slightly in shape from high-vacuum welds). Weld shrinkage across nonvacuum electron-beam welds is generally less than half that of gas-tungsten arc welds on the

same work metal and thickness and is typically about twice that for high-vacuum electron-beam welds. Tooling generally is designed specifically for welding a particular assembly in mass production. Providing fixtures and

equipment that permit welding at high speeds and efficient work handling is the key to obtaining high production rates and low cost. The work-handling equipment is at atmospheric pressure instead of in a vacuum; hence, general-purpose welding positioners and related fixtures usually are satisfactory. Hold-down clamping devices can be simpler than those often needed for arc welding similar parts, because angular distortion or dihedral warpage is generally low. Locating and alignment mechanisms are simpler than for welding in a vacuum, because of their accessibility and because beams are

wider and lower accuracy is required in tracking the joint. Maintenance of work-handling equipment is greatly simplified by its out-of-chamber location. Applications. Nonvacuum electron-beam welding is used in the production of transmission-train components, steering

column jackets, die-cast aluminum intake manifolds, and catalytic converters. In addition, several appliance (and other consumer-related) items also are produced using the EBW-NV process. Figure 19 shows a catalytic converter that was nonvacuum electron-beam welded with the beam in a horizontal position.

FIG. 19 CATALYTIC CONVERTER ASSEMBLY CONSISTING OF SEAMS JOINED BY EBW-NV PROCESS


Controlling Heat Effects The extent of heat-affected base metal in EBW is controlled mainly by regulating the welding speed, width, and shape of the weld. In general, the width of the HAZ is proportional to the average width of the weld, which also strongly influences shrinkage and distortion. Damage to heat-sensitive attachments, inserts, or encapsulated materials close to the weld joint is controlled mainly by regulating the width and shape of the weld bead. Damage to nearby finished functional surfaces on the workpieces or to nearby drilled or tapped holes or other crucial features of the workpieces can often be avoided in EBW because of the narrowness of the welds. Joint designs that provide constricted sections near the weld to retard heat transfer from the weld, or that provide an integral heat sink, and the use of external heat sinks or chill bars are also helpful in controlling heat effects in EBW. Narrow, closely controlled electron-beam welds, with low heat input, can be made in close proximity to heat-sensitive attachments, inserts, or encapsulated materials without damaging them, often when no other welding process can meet this requirement. Procedure Development and Practice Considerations for Electron-Beam Welding

Welding of Thin Metal Electron-beam welding can be advantageous for joints in which at least one member is made of thin metal (610 mm/min, or 24 in./min) selected to minimize heat input can promote pronounced directionality of grain growth in the weld pool, leading to a plane of weakness at the weld centerline. High-frequency (>1

kHz) beam oscillation can be applied to restore an elliptical weld-pool geometry at high travel speeds and eliminate the pronounced delineation of the weld centerline. As with any low-ductility weldment, the degree of restraint imposed on the joint during the welding operation should be minimized. Most refractory metals and alloys can be welded in thin (5.1

100500 >200

~1.5

~60

COMMENTS CONCERN FOR WHISKER GROWTH WITH MOISTURE AND RESIDUAL STRESSES; RECOMMENDED FUSING MELTING POINT OF TIN, 232 °C (450 °F) AVAILABLE COMPOSITIONS: 63SN-37PB, 95PB5SN, ETC.; PLATING COMPOSITION MAY DEVIATE FROM NOMINAL VALUES MELTING POINTS: 63SN-37PB, 183 °C (361 °F); 95PB-5SN, 314 °C (597 °F) UNIFORM COVERAGE REQUIRED PROTECTION SENSITIVE TO DEPOSITION PROCESS DETAILS; QUALIFICATION TESTS RECOMMENDED

Source: Ref 24

For metals such as aluminum and magnesium, which simply cannot be kept solderable, or for ceramics that are intrinsically unsolderable, a metal layer is deposited on the substrate. The solder more readily wets this layer, which is often referred as to the "solderable" layer. These films must be adequately thick to avoid being consumed by the metallurgical reaction with the solder, and they must exhibit adequate adhesion to the base material. In some instances, an "adhesive" layer may be deposited between the base material and the solderable layer. A "protective" layer can also be deposited onto the solderable coating in order to prevent excessive oxidation or contamination. This protective layer is then consumed by the solder. Table 36 lists the coating materials and processing steps used for substrates that are commonly involved in solder processing.

TABLE 36 SOLDERABLE AND PROTECTIVE FINISHES FOR COMMON SUBSTRATE MATERIALS

SUBSTRATE

SOLDERABLE LAYER

COPPER

ELECTROPLATED NICKEL(A) ELECTROPLATED COPPER

COPPER ALLOYS WITH ZINC NICKEL AND NICKEL ALLOYS

ALUMINUM AND ALUMINUM ALLOYS(D)

IRON,IRONBASE ALLOYS (ALLOY 42, ALLOY 52, KOVAR), LOW-CARBON STEELS, AND STAINLESS STEELS

CAST IRON(D)

CERAMICS (A12O3, BEO, SIO2, SI3N4, ETC.)(D)

COMMENTS

SOLDERABLE LAYER THICKNESS μM μIN. 1.560-150 3.8 2.5(C) 100(C)

PROTECTIVE LAYER

ELECTROPLATED GOLD(B) ...

PROTECTIVE LAYER THICKNESS μM μIN. 1.350-100 2.5 ... ...

ELECTROPLATED NICKEL(A) ELECTROPLATED SILVER

1.53.8 3.88.9

ELECTROPLATED GOLD(B) ...

1.32.5 ...

...

ELECTROPLATED COPPER ELECTROPLATED COPPER (ZINCATE UNDERCOAT)

3.87.6 3.8-25

150300 3001000

...

...

...

...

...

...

ELECTROPLATED NICKEL (ZINCATE UNDERCOAT)(A) ELECTROPLATED NICKEL(A)

1.33.8

50-150

ELECTROPLATED GOLD(B)

1.32.5

50-100

1.33.8

50-150


1.32.5

50-100

ELECTROPLATED SILVER ELECTROPLATED PALLADIUMNICKEL ELECTROPLATED IRON

3.88.9 1.33.5

150350 50-150

...

...

...

FOR COATED STEELS, USE RECOMMENDED SOLDERABLE AND PROTECTION FINISHES FOR THE PARTICULAR COATING MATERIAL ...

...

...

...

...

1.33.8

50-150

...

...

...

PRECIOUS METAL THICK FILMS

10-20

100800


2.02.5

80-100

REFRACTORY THICK FILMS MAY INCLUDE NICKEL

10-20

100800 40-350

...

...

...

EXTENDED THICKNESS TO 2.5 TO 7.6 μM (100 TO 300 μIN.) FOR IMPROVED COVERAGE CERAMIC SURFACE ROUGHNESS AND THICK-FILM POROSITY MAY REQUIRE THICKER GOLD DEPOSITS ...

...

...

...

...

1.09.0

60-150 150350

50-100

... 12 μM (5000 μIN.) FOR ELEVATEDTEMPERATURE SERVICE ... SMALL SOLDERABILITY LOSS THAT IS DUE TO SILVER TARNISH ... ANODIC FINISH (CD OR ZN) RECOMMENDED TO PREVENT CORROSION BETWEEN ALUMINUM AND COPPER OR NICKEL ...

(SOLDERABLE OR BARRIER LAYER BARRIER LAYER)

(A) THICKER LAYERS OF 2.5 TO 7.6 μM (100 TO 300 μIN.) PERMITTED FOR IMPROVED COVERAGE; ELECTROLESS NICKEL, 1.3 TO 3.0 μM (50 TO 120 μIN.) OR 2.5 TO 5.1 μM (100 TO 200 μIN.) PERMITTED ON NONFLEXIBLE SUBSTRATES; LAYER ADHESION AND SOLDERABILITY REQUIRE CLOSE MONITORING. (B) THICKER LAYERS OF 2.5 TO 7.5 μM (100 TO 300 μIN.) PERMITTED FOR SOLDERABLE LAYER PROTECTION; HOWEVER, THINNER LAYERS ARE RECOMMENDED AND SOLDER JOINTS SHOULD BE MONITORED FOR EMBRITTLEMENT. (C) MINIMUM. (D) NOT CONTAINED IN MIL-STD-1276D Gold plating layers should be removed by hot-solder dipping twice in a bath of molten solder if the calculated gold content of the solder joint will exceed 3 to 4 wt%, in order to prevent solder-joint embrittlement by gold-tin intermetallic formation. Electroplated gold layers should be pure gold, preferably type III, 99.9% pure, and grade A, according to MILG-45204C. Gold coatings alloyed with cobalt, nickel, or both (termed "hard gold"), which are used for wear resistance, are difficult to solder, because of oxidation of the cobalt or nickel component. Only matte finishes should be specified for soldering applications. "Bright" platings require organic additives in the bath, which then become entrapped in the gold coating and subsequently volatilize during soldering. This creates voids in the joints and poor wetting properties.


21. R. KLEIN-WASSINK, SOLDERING IN ELECTRONICS, ELECTROCHEMICAL PUB. LTD., AYR, SCOTLAND, 1989, P 189 22. SOLDERING MANUAL, AWS, 1977, P 35-39 23. P. VIANCO ET AL., SOLDERABILITY TESTING OF KOVAR WITH 60SN-40PB SOLDER AND ORGANIC FLUXES, WELD. J. RES. SUPPL., JUNE 1990, P 230S 24. BRAZING HANDBOOK, AWS, 1991, P 267-276 General Soldering Paul T. Vianco, Sandia National Laboratories

Fluxes The major role of the flux is the removal of thin tarnish layers during the initial stages of the soldering process, thereby permitting the molten solder to react with the substrate and to spread. Fluxes do not effectively displace organic contaminants, so the substrate underneath such films will remain untouched by the flux and will therefore be unsolderable. A degreasing step should precede flux application. Even the strongest fluxes cannot remove thick oxides or heavy scales, so a mechanical or chemical precleaning procedure may be required. The flux has two additional functions. One is that it lowers the surface tension of the solder, allowing it to more readily fill gaps and holes by capillary action. The other function is that the flux coating protects the metal surface from reoxidation during the heating steps just prior to soldering. Fluxes contain three principal ingredients: an active chemical compound, such as a halide, for oxide removal; wetting agents to improve surface coverage; and a vehicle to dilute and mix the cleaning compound and wetting agents together. The vehicle, which is removed by evaporation during the soldering process, is typically water, isopropyl alcohol, glycerin, glycol (for liquid fluxes), or petroleum jelly (for flux pastes or creams). Fluxes are characterized by their cleaning agent

and are assigned to one of these categories of increasing activity: rosin-base fluxes, organic-acid fluxes (also called "intermediate" or "water-soluble" fluxes), and inorganic-acid fluxes. Flux specifications are most numerous for the rosin-base materials, because of the criticality of their properties in the assembly of electronic products. Sample specifications include military standards MIL-F-14256E and MIL-STD-2000; federal specification QQ-S-571 E; and the industry standards ASTM D 509, IPC-S-815A, and IPC-TM-650 (test methods). Fluxes are tested for bulk corrosivity (copper mirror test), concentrations of Cl- and F- (halide content), and solids content. Tests of residues left by the flux on the substrate (typically a printed wiring board) include surface insulation resistance, and ionic residues. The ISO/DIS 9454-1 document includes organic- and inorganic-acid flux specifications and test procedures (ISO/DIS 9455-X, where X = 1, 2, . . ., 14), in addition to rosin-base flux specifications. The removal of flux residues depends on the specific application, as well as the potential corrosivity of the residues left behind. If residues are specified for removal, then the workpiece should be cleaned as soon as possible after processing. Rosin residues become particularly difficult to remove with time. Rosin-base fluxes contain "water-white" rosin, a distillation product from pine tree sap. When used alone, rosin-base

fluxes are referred to as a "nonactivated," or type R, grade. Rosin is solid, noncorrosive, and insulating at room temperature. Heating causes it to liquefy and become slightly chemically active, because of its abietic and pimaric acid components. The rosin is typically dissolved in an organic vehicle, such as isopropyl alcohol, to form the liquid solution. The addition of an activator to rosin fluxes increases their chemical activity. Activators can be organic halogenated compounds, such as amine hydrohalides that contain chloride, fluoride, or bromide ion groups or "halide-free" activators, such as oleic, stearic, or lactic acids. Halide-free fluxes are recommended for materials that are sensitive to stresscorrosion cracking. The concentration of activators, which defines the corrosivity of the flux, determines the flux category as being one of the following: rosin-base, mildly activated (RMA), fully activated (RA), and superactivated (SA). Generally, rosin-base fluxes have poor high-temperature stability; that is, the fluxes degrade rapidly after excessive exposure to elevated temperatures (for example, >25 s at 260 °C, or 500 °F). R-type fluxes are used on electronic assemblies, where the benign residues are left on the soldered assembly. However, the substrate(s) must have excellent solderability, because the weak activity of these fluxes offers a very small window of variability for oxide thicknesses. R-type fluxes are effective on lightly tarnished copper surfaces and on precious-metal substrates, as well as on tin- and solder-coated base metals. Residues can be removed by nonpolar solvents, organic solvents (isopropyl alcohol, 1,1,1-trichloroethane), semi-aqueous solutions (terpenes), and aqueous baths or detergent solutions. The residues left by R-type fluxes can be reduced by dilution with an organic solvent (for example, isopropyl alcohol). Such fluxes are termed "low solids" because the solids, which are responsible for the residues, are significantly reduced so that cleaning steps are not required. These fluxes are also referred to as "no-clean;" however, the term "no-clean" covers a much wider range of flux formulations that must simply comply with residue corrosivity, irrespective of the solids content. The increased activity of the RMA fluxes improves solder wetting on more heavily tarnished copper substrates. Residues are relatively benign and only require removal when the fluxes are used on high-reliability electronic systems (for example, in military applications) or when the fluxes contain potentially harmful halide activators. Residue cleaning is a two-step process. The first step requires the use of organic solvents to remove the nonpolar rosin, and in the second step, a polar solvent (water or polar alcohols) is used to rinse away activator residues. Aqueous and semi-aqueous products can perform both functions. The greater corrosivity of RA and SA fluxes is due, primarily, to higher levels of the activators that are used in RMA fluxes. These fluxes are used on base metals such as nickel (and nickel plate), lightly tarnished low-carbon steels, copper, iron-base alloys, and copper-base alloys (brasses, bronzes, and beryllium copper). Residues of the RA and SA fluxes must be removed to prevent the corrosion of the substrate later in service. Cleaning procedures similar to those identified for RMA fluxes are used. Cleanliness testing for ionic residues should be implemented to verify the adequate removal of flux deposits, particularly for the stronger fluxes. Organic-acid fluxes have a chemical activity range that varies from the levels associated with RMA fluxes to those that

exceed the activity of RA fluxes. These fluxes contain one or more organic acids, such as lactic, oleic, or stearic acid.

Chemical activity can be enhanced by adding organic halogen compounds (amine hydrohalide, which may contain chloride and bromide derivatives) or nonhalogenated substances, such as one of the amines or amides (urea or thylene diamine). Typical vehicles include water, isopropyl alcohol, polyglycols, or petroleum jelly for pastes. The organic-acid fluxes are used in many electronic applications involving machine processes and hand assembly, as well as the hot tin or solder dipping of nickel and iron-base alloy leads and devices. These fluxes are also used on structural applications with copper and copper alloy workpieces that have light to moderate tarnishes. Whether the fluxes are used in structural or electronic applications, their residues should be removed. The water-soluble residues are typically removed by rinsing with water or a polar organic solvent. The organic acid fluxes have better hightemperature stability, when compared with the rosin-base materials, which makes them good candidates for solders with melting temperatures that exceed 200 °C (363 °F). Inorganic-acid fluxes have the highest levels of chemical activity. There are two categories of these fluxes: • •

PURE ACIDS, SUCH AS HYDROCHLORIC, HYDROFLUORIC, OR PHOSPHORIC ACIDS, WHICH HAVE SURFACTANTS ADDED TO ENHANCE COVERAGE. INORGANIC SALT MIXTURES OR SOLUTIONS, WHICH MAY ALSO CONTAIN SURFACTANTS.

The pure acids are very strong and are capable of removing heavy oxide layers and scales. However, they offer the clean surface limited protection from reoxidation during the soldering process. Binary inorganic salt "alloys" are formed from the combination of zinc chloride (ZnCl2), ammonium chloride (NH4Cl), stannous chloride (SnCl2), or sodium chloride (NaCl) components. The inorganic salt combinations form simple eutectics with minimum temperatures near the melting point of solders. (The melting temperature of the single salts are too high for them to flow at typical soldering temperatures.) For example, the NH4Cl:3 (ZnCl2) eutectic composition has a eutectic temperature (Te) of 177 °C (351 °F). Other eutectic salt combinations include: SnCl2-CuCl, Te = 170 °C (338 °F); AgCl-CuCl, Te = 260 °C (500 °F); and CuClZnCl2, Te = 230 °C (446 °F).

Proper flux selection is based on the salt mixture having a melting temperature that is less than that of the solder. The molten salt reduces the surface oxides and coats the base metal to prevent reoxidation during the soldering process. Flux activity is further increased by dissolving the flux in a water vehicle. The salts break down to release Cl- ions, which combine with water to form a hydrochloric acid (HCl) "activator." The fluxes, particularly the water solutions, may contain surfactants to assist the complete coverage of the joint area. A paste form of the flux uses petroleum jelly as the vehicle. The inorganic-acid fluxes are limited to structural applications, such as plumbing or mechanical assemblies. Their corrosive activity is unacceptable for electronic devices or substrates, or for their assembly. These fluxes are effective on nickel and nickel alloys, stainless steels, chromium, and heavily tarnished copper and copper alloys. The inorganic fluxes are very stable at high temperatures. The flux residues and the fluxes themselves are extremely corrosive and must be thoroughly removed after processing. It should be noted that chloride-containing residues and small amounts of water are particularly damaging to aluminum and aluminum alloys, because of the potential for stress-corrosion cracking. Residues are removed with hot-water rinses and follow-up polar-solvent rinses. Neutralizing solutions (mild caustics) can be used on nonsensitive substrates. A list of selected inorganic flux solutions for particular base metals is provided in Table 37.

TABLE 37 TYPICAL INORGANIC FLUXES FOR SELECTED BASE METALS

FOR STAINLESS STEEL AND GALVANIZED IRON ZINC CHLORIDE, ML (OZ) 2510 (85) AMMONIUM CHLORIDE, ML (OZ) 190 (6.5) STANNOUS CHLORIDE, ML (OZ) 270 (9) HYDROCHLORIC ACID, ML (OZ) 60 (2) WETTING AGENT (OPTIONAL), WT% 0.1 (A) WATER FOR STAINLESS STEEL ZINC CHLORIDE, ML (OZ) 1420 (48) AMMONIUM CHLORIDE, ML (OZ) 150 (5) HYDROCHLORIC ACID, ML (OZ) 90 (3) WETTING AGENT (OPTIONAL), WT% 0.1 WATER(A) FOR MONEL ZINC CHLORIDE, ML (OZ) 470 (16) AMMONIUM CHLORIDE, ML (OZ) 470 (16) GLYCERIN, ML (OZ) 470 (16) WATER, L (GAL) 0.5 (0.125) FOR HIGH-TENSILE-STRENGTH MANGANESE, BRONZE, COPPER, OR BRASS ORTHOPHOSPHORIC ACID (85%), ML (OZ) 1000 (34) WATER, ML (OZ) 470 (16) FOR CAST IRON ZINC CHLORIDE, ML (OZ) 950 (32) AMMONIUM CHLORIDE, ML (OZ) 120 (4) SODIUM CHLORIDE, ML (OZ) 240 (8) HYDROCHLORIC ACID, ML (OZ) 240 (8) WATER(A) FOR CAST IRON ZINC CHLORIDE, ML (OZ) 1180 (40) AMMONIUM CHLORIDE, ML (OZ) 120 (4) HYDROFLUORIC ACID, ML (OZ) 40 (1.25) WATER(A) PASTE FLUX FOR SOLDERING ALUMINUM STANNOUS CHLORIDE, ML (OZ) 2450 (83) ZINC DIHYDRAZINIUM CHLORIDE, ML (OZ) 210 (7) HYDRAZINE HYDROBROMIDE, ML (OZ) 295 (10) WATER, ML (OZ) 295 (10) FOR SOLDERING ALUMINUM CADMIUM FLUOBORIDE, ML (OZ) 150 (5) ZINC FLUOBORIDE, ML (OZ) 150 (5) FLUOBORIC ACID, ML (OZ) 180 (6) DIETHANOL AMINE, ML (OZ) 590 (20) DIETHANOL DIAMINE, ML (OZ) 120 (4) DIETHANOL TRIAMINE, ML (OZ) 295 (10) POTASSIUM CHLORIDE, ML (OZ) 1330 (45) SODIUM CHLORIDE, ML (OZ) 890 (30) LITHIUM CHLORIDE, ML (OZ) 440 (15) POTASSIUM FLUORIDE, ML (OZ) 210 (7)

SODIUM PYROPHOSPHATE, ML (OZ) TRIETHANOLAMINE, ML (OZ) FLUOBORIC ACID, ML (OZ) CADMIUM FLUOBORATE, ML (OZ) STANNOUS CHLORIDE, ML (OZ) AMMONIUM CHLORIDE, ML (OZ) SODIUM FLUORIDE, ML (OZ) ZINC CHLORIDE, ML (OZ) AMMONIUM CHLORIDE, ML (OZ) SODIUM FLUORIDE, ML (OZ) (A)

90 (3) 740 (25) 90 (3) 60 (2) 1300 (44) 150 (5) 30 (1) 1300 (44) 150 (5) 30 (1)

TO MAKE 3.8 L (1 GAL)

General Soldering Paul T. Vianco, Sandia National Laboratories

Solder Joint Assembly The configuration of a solder joint must have two basic attributes: It must meet the service requirements of the assembly, and it must be manufacturable to those requirements. Specifications and guidelines for the design of structural solder joints (for example, conduit) are limited, when compared with those used in electronic applications. Some British and European standards have been established, but few exist in the United States. Because standard joint design guidelines are not available, performance data are not well documented. However, the brazing community has several sets of standards that pertain to structural brazed joints (Ref 24). These specifications, which were established by the American Society of Mechanical Engineers (ASME), American National Standards Institute (ANSI), American Welding Society (AWS), and the Society of Automotive Engineers (SAE), cover joint construction, strength tests, and safety issues. Although these specific documents were developed for brazements, they can provide general guidelines for applying similar tests and criteria on solder joints. Under the circumstances, conservative design practices and inspection and test protocols that certify process quality have substituted for insufficient data. Additional information is available in the article "Evaluation and Quality Control of Soldered Joints" in this Volume. Design for Service Requirements. Solder joint service conditions are typically represented by either electronic interconnect applications or structural applications. The former are described in the article "Soldering in Electronic Applications." Structural applications, which are described below, range from the assembly of small housings and containers to plumbing for fluid or gas conduit. Typical service requirements for structural joints include the support of mechanical loads (both high-temperature and cryogenic conditions), hermetic sealing for vacuum applications, and leaktight joints for positive-pressure conduits and containers.

Several common structural solder joint configurations are shown in Fig. 15. It is recommended that solder joints be loaded in shear. Tensile loads should be avoided, because a slight deviation in the loading direction can place the joint in a peel-type fracture mode, which typically causes premature failure of the assembly. Lap-shear joints should have an overlap of approximately one-third of the base-metal thickness (Ref 25). Lap joints retain their strength over a range of gap thicknesses. However, the gap thicknesses of a single joint should be uniform for consistent capillary filling by the molten solder.

FIG. 15 COMMON STRUCTURAL SOLDER JOINTS

One of the first of the design criteria that pertains to the selection of a suitable solder is melting temperature, specifically the solidus and liquidus points. The solidus temperature is determined directly by the service conditions of the final assembly. Of course, the service temperature should never exceed the solidus point. However, how close the service temperature should be allowed to approach the solidus temperature depends on the level of monotonic or creep deformation and the fatigue life acceptable for the particular application. Tables of appropriate data either should be referred to or the necessary experiments performed. A suitable liquidus temperature is primarily determined by the temperature sensitivity of the substrate materials during the soldering process. The molten solder working temperature should be approximately 25 to 50 °C (45 to 90 °F) greater than the liquidus temperature. A secondary factor in the selection of the liquidus and solidus temperatures is the desired pasty range for solder workability prior to solidification. Finally, multiple soldering processes require the solidus temperature of the first (high-temperature) soldering step to exceed the liquidus point by 30 to 50 °C (55 to 90 °F) of the second (lowtemperature) step. The mechanical properties of the solder must be considered in terms of the particular application. For example, the monotonic and creep strengths (shear or tensile) must be appropriate for the particular loading conditions and temperature environments to which the joint is subjected. Fatigue resistance is important for joints subjected to time-varying loads. Although such loads may be well below the ultimate monotonic or creep strength limits of the solder, they can result in significant cracking of the solder layer in a joint. The size of the substrates and the thermal conductivity of the material from which they are constructed affect the joining process. A disparity of mass between the two substrates will result in the larger part heating up more slowly than the smaller one. Insufficient heat can result in a cold joint (that is, an absence of capillary flow and/or poor fillet formation

caused by premature solder solidification). Conversely, overheating the smaller workpiece may cause undesirable changes to the mechanical and physical properties of the base material. Similarly, two substrates with vastly different thermal conductivities (for example, stainless steel and aluminum) are difficult to uniformly heat, particularly in localized heating processes, such as soldering with an iron or torch. When designing joints of dissimilar materials, the coefficient of thermal expansion (CTE) differences must be considered because sufficient stresses can build within the joint to cause catastrophic failure, particularly if the substrates are brittle. Time variations in environmental temperatures cause fatigue deformation to accumulate in the solder, thereby weakening the mechanical strength of the joint. An example of the method used to estimate the fatigue life of a 60Sn-40Pb solder joint can be made by calculations based on empirical data, as shown below. Unfortunately, data to support similar computations with other solders have not been compiled. Laboratory tests of the isothermal fatigue life of 60Sn-40Pb over the temperature range from -50 to 125 °C (-58 to 257 °F) were used to develop the following equation for the fatigue life, Nf (Ref 16): 1.96

 1.14  Nf =   ∆γ   p

where Nf is the number of cycles to failure (50% load drop for equivalent strain) and ∆γp is the plastic strain level in the joint. A conservative estimate of ∆γp is made by using the total strain, ∆γt, determined by: ∆γ t =

(α1 − α 2 )∆TL t

where αi is the CTE of the two substrates (i = 1, 2), ∆Tis the temperature difference, L is the maximum dimension of the joint, and t is the joint thickness. The parameters and joint configuration are illustrated in Fig. 16. Solder fatigue cracks may be visible in the solder fillets, typically where L is a maximum.

FIG. 16 PARAMETERS FOR THE CALCULATION OF TOTAL THERMAL MISMATCH STRAIN IN A GENERIC SOLDER JOINT CONFIGURATION. T, JOINT THICKNESS; L, MAXIMUM JOINT DIMENSION BETWEEN THE TWO SUBSTRATES

Thermal expansion mismatch between workpieces affects joint clearances during the soldering process. Consider the enclosed geometry tube-in-socket joint in Fig. 17. When the CTE of the outer part is larger than that of the inner part, the gap will increase with heating. An excessive gap will result in voids in the joint that are caused by an insufficient solder supply and/or poor capillary action. A sample calculation follows: • •

INNER PART AT 20 °C (68 °F), COPPER; CTE (α1), 17.0 × 10-6/K; OUTER DIAMETER, 74.7 MM (2.94 IN.) OUTER PART AT 20 °C (68 °F), ALUMINUM; CTE (α0), 23.5 × 10-6/K; INNER DIAMETER, 75.0

•

MM (2.95 IN.) GAP AT 20 °C (68 °F), 0.15 MM (0.006 IN.)

FIG. 17 EFFECTS OF HEATING ON THE JOINT GAP BETWEEN DISSIMILAR METALS

At 250 °C (482 °F), the new dimension, Lf, of the parts (the outer diameter of the aluminum and the inner diameter of the copper) are calculated by:

LF = LO + LOα∆T where Lo is the dimension at 20 °C (68 °F), αis the CTE of the material, and ∆T is the temperature difference (Tf - To). • • • •

TF = 250 °C (482 °F), SO THAT ∆T = 230 °C (414 °F) INNER PART AT 250 °C (482 °F), COPPER; OUTER DIAMETER, 75.0 MM (2.95 IN.) OUTER PART AT 250 °C (482 °F), ALUMINUM; INNER DIAMETER, 75.4 MM (2.97 IN.) GAP AT 250 °C (482 °F), 0.2 MM (0.008 IN.)

In this example, the solder joint gap grew by 0.05 mm (0.002 in.). This computation can be applied to any geometry in which Lo is the dimension of the respective parts. Conversely, an inner part with a CTE that is greater than that of the outer section will cause the gap to decrease with heating, causing voids and poor solder filling resulting from the inability of flux volatiles to escape from the joint. Finally, consideration must be given to material compatibility. The first concern is that the substrate materials are not damaged by chemicals or the mechanical procedures used to preclean surfaces. Damage includes excessive metal removal, stress-corrosion cracking, or the deposition of hard-to-remove residues. Secondly, the corrosion potential of joining dissimilar metals, including the two base metals as well as the solder/base metal couples, must be adequately assessed. In the case of conduit for liquids or gases, the role of the carried medium, as an electrolyte in the corrosion of the joint, must also be considered. In many cases, references to the electromotive series or fundamental chemistry

textbooks (Ref 26) can provide corrosion information for a preliminary assessment. Finally, the assembled product must be compatible with the chemicals and processes used to remove flux residues after soldering. Design for Processing Requirements. It is assumed that the substrates have adequate solderability for the selected solder alloy. Two processes must occur during product assembly to ensure an adequate joint: The solder must fill the gap area and flux volatiles and liquids must be ejected from the joint. Some general guidelines at the design stage will facilitate these processes.

First, the joint gaps should be in the range from 0.076 to 0.18 mm (0.003 to 0.007 in.). Workpiece design must also account for changes in gap dimensions caused by the heating/cooling of dissimilar substrates, surface roughness, plate flatness, eccentricity of circular-section workpieces, and consumption of the base metal or substrate coatings by the molten solder. Smaller gaps restrict solder flow and hinder the removal of flux volatiles, whereas larger gaps diminish the capillary effect required by the molten solder to fill the joint. Second, the preferred gap geometry is a straight path opened to the atmosphere at two locations. One opening is used to supply flux and solder to the joint, while the other serves to vent flux volatiles and gases as the molten solder fills the gap. "Blind" gaps (or blind holes), which do not allow a vent path, should be avoided. A vent hole can be added and later filled in with a lower-melting-temperature solder if the design permits. Another rule of thumb is that solder has difficulty in "turning corners." Therefore, straight paths are preferred over sharp comers in the joint. Third, ready access to the joint by a reservoir of solder containing an adequate quantity to fill the gap and to form the fillet must be provided. Inadequately filled gaps can be caused by an insufficient supply of solder to the joint. Poorly formed solder fillets reduce the monotonic strength, fatigue life, and creep strength of the joints, as well as limit the solderability assessment. Fourth, solder joints that use preforms are always susceptible to void formation. Because the joint does not fill by capillary action, volatiles and gases have an increased likelihood of being trapped within the gap. Such joints must also be carefully assessed for corrosion potential by retained flux residues in those voids. Fifth, the metallurgical reaction between the substrate and the solder during wetting necessarily consumes some of the substrate to form the reaction layer (typically, an intermetallic compound) at the interface. The dissolution of substrate material must be accounted for when soldering to very thin foils or coatings. The rate of dissolution will depend on substrate geometry (round or cylindrical sections dissolve slightly faster than flat plates), substrate composition, and the composition and volume of solder (flowing baths will remove more substrate metal than static baths). The temperature of the solder and the time of immersion (an increase of both raise the rate of substrate dissolution), are also factors. Dissolution data from unspecified geometries are determined from Fig. 18 and 19.

FIG. 18 NOMOGRAPH TO DETERMINE THE TIN-LEAD SOLDER REACTION ZONE THICKNESS FOR THE COPPER ALLOYS C10200 AND C26000 ALLOYS UNDER STATIC INFINITE-VOLUME CONDITIONS. SELECT THE EXPOSURE CONDITIONS ON SCALES A AND B. WITH A STRAIGHT EDGE BETWEEN THE CONDITIONS, FIND THE INTERSECTION WITH BASE LINE C. USING A STRAIGHT EDGE, CONNECT THIS POINT ON LINE C WITH THE

TIN CONTENT OF THE SOLDER ON SCALE D. READ THE REACTION-ZONE THICKNESS ON SCALE E. THE NOMOGRAPH CAN ALSO BE USED IN REVERSED SEQUENCE TO SELECT OPERATING CONDITIONS, GIVEN A PERMISSIBLE REACTION ZONE THICKNESS AND SOLDER COMPOSITION.

FIG. 19 METAL DISSOLUTION RATES FOR ELEMENTAL METALS AS A FUNCTION OF TEMPERATURE. SOURCE: REF 27

Intermetallic compounds represent the metallurgical reaction product between the solder and the substrate material. The compound formulations for common base metals and coatings with tin-bearing solders are: • • • • •

COPPER: CU6SN5 AND CU3SN NICKEL: NI3SN4 GOLD: AUSN4 AND AUSN2 IRON: FESN2 SILVER: AG3SN

These layers form when the solder is molten, as well as after solidification, by means of solid-state diffusion during storage or service at elevated temperature. In the solid state, the thickness is a function of the temperature and the time. Although intermetallic compounds are a necessary part of solder wetting, excessively thick layers can be detrimental to joint integrity or follow-up processing. Specifically, the brittle nature of intermetallic compounds can jeopardize the mechanical and fatigue strength of the joint, particularly at high loading rates, such as those that occur under impact or vibration conditions. Another detriment is that the solid-state growth of the layers consumes both the solder film and the base metal. Complete consumption of the tin or solder coating exposes the intermetallic surface to oxidation, making it very difficult to solder coat, even when using the strongest fluxes. The solid-state growth kinetics of several solders that were hot dip coated on copper are shown in Fig. 20. The typical as-soldered thickness of intermetallic compounds of

60Sn-40Pb on copper and nickel are 1 to 2 μm (39 μin.) and 0.5 μm (20 μin.), respectively. The layers grow very slowly at room temperature.

FIG. 20 SOLID-STATE INTERMETALLIC GROWTH KINETICS FOR HOT-DIPPED SOLDER COATINGS ON COPPER. SOURCE: SANDIA NATIONAL LABORATORIES

Joint design must include postassembly cleaning requirements. Materials compatibility between the cleaning agents and the workpiece, as well as the geometry of the workpiece, are important factors. Reentrant comers or hidden gaps and passages can entrap flux residues and prevent the flow of clean solvents into the areas. The use of the high-temperature, precious-metal solders 80Au-20Sn or 88Au-12Ge on precious-metal substrates and coatings does not require a flux to remove the very limited oxide layers. The melting points of these solders preclude the use of rosin-base fluxes, because of significant high-temperature degradation. Other fluxes may also experience some degradation. The absence of a flux causes the surface tension of the solder to remain high, impeding capillary flow by the molten alloy. Therefore, the solders are typically introduced into the joint as preforms. Mechanical pressure on the joint is used to assist the spreading of the solder throughout the gap. Small quantities of oxide particles in the solder arise from the non-precious-metal component. Random spreading of the solder causes the joint to have a higher number of voids than is observed in a joint filled by capillary action. Maintaining the position of the substrates during solder wetting is necessary for successful joint formation. Special fixturing can be designed to hold the workpieces for critical applications. Another approach is to design the joint itself to be self-jigging, in order to maintain the proper gaps and clearances during soldering (Fig. 21).

FIG. 21 SELF-JIGGING JOINT CONFIGURATIONS

Design for Inspection Requirements. The inspection techniques selected to qualify a solder joint will have particular

requirements in terms of joint configuration. Visual inspection is typically used to determine whether the solder has flowed within the gap and on the external surfaces (fillets). For blind holes or gaps, witness holes can be drilled into the substrate(s). During soldering, the solder fills the witness hole to the outside surface, indicating to the operator that the solder has indeed filled the gap. Although they are not always practical, unobstructed views of the solder joint gap at the external surface vastly improve the visual inspection of the joint.


16. H. SOLOMON, "FATIGUE OF 60/40 SOLDER," REPORT 86CRD024, GENERAL ELECTRIC CO., MAY 1986 24. BRAZING HANDBOOK, AWS, 1991, P 267-276 25. C. THWAITES, CAPILLARY JOINING, WILEY AND SONS, UNITED KINGDOM, 1982, P 52 26. J. SCULLY, THE FUNDAMENTALS OF CORROSION, 2ND ED., PERGAMON PRESS, INC., 1975 27. C. LEA, A SCIENTIFIC GUIDE TO SURFACE-MOUNT TECHNOLOGY, ELECTROCHEM. PUBL., LTD., 1988 General Soldering Paul T. Vianco, Sandia National Laboratories

Soldering Techniques The numerous soldering processes have the common goal of providing heat energy to the joint area in the presence of the flux and solder alloy. The heat energy (and subsequent temperature rise) must support activation of the flux cleaning activities, heating of the substrate, and melting of the solder alloy to permit wetting and joint formation. Too little heat limits solder wetting, whereas excessive heat chemically degrades the flux (which loses activity and deposits hard-toremove residues), reoxidizes the substrate surface prior to completion of solder wetting, or causes heat damage to the substrate. Heating techniques can be categorized as global processes, in which the entire substrate reaches the soldering temperature, and local techniques, in which only the immediate joint area is heated. Examples of global techniques include vapor-phase (condensation) reflow, furnace (infrared or convection) reflow, wave soldering, ultrasonic soldering (bath), as well as dip (including hot-air leveling) and drag soldering. Local heating techniques include the soldering iron (heat and heat plus ultrasonics), torch, hot gas reflow, focused infrared soldering, induction soldering, and laser soldering. Operator interaction depends on the particular process. For example, the entire process can be fully automated, as in furnace systems or specialized equipment used in electronic circuit board assembly. Nevertheless, operators must establish the proper zone temperatures, conveyor belt speeds, and heating rates. Thermocouple profiling of a prototype workpiece is essential to verifying that the proper thermal conditions will be achieved. Manual soldering, such as with the iron or torch, requires greater operator interaction and skill. Control of the full sequence of heating, fluxing, and applying the filler metal is particularly critical in manual solder processes (whether iron, induction, torch, or hot-gas processes), because they are usually performed in air. Competing factors that occur during substrate heating are: • • • •

NEED FOR AN INCREASE IN SUBSTRATE TEMPERATURE TO MAINTAIN THE MOLTEN STATE OF THE SOLDER ACTIVATION OF THE FLUX OXIDE REMOVAL PROCESS (OR PROCESSES) ACCELERATED OXIDATION OF THE SUBSTRATE SURFACE DEGRADATION OF SOME FLUX COATINGS

Circuit board soldering and structural applications are described below as general examples. Circuit Board Soldering. First, flux is applied to the leads and circuit board, unless flux-cored soldering wire is used, in

which case this step is omitted. The coated area should extend beyond the immediate joint to ensure adequate wetting by the solder. Next, the soldering iron is made to contact the lead (never the circuit board land), and the solder wire is fed so that it contacts the opposite side of the lead. Melting of the solder indicates that the lead has reached temperature. The wire, followed by the iron tip, is then removed from the joint area upon formation of the joint fillet. Structural Applications. The larger thermal mass of typical workpieces requires preheating in order to bring the joint

area to the working temperature. Then, the flux is applied to fully cover the hot surface. The joint is further heated and the filler metal is applied to the joint surfaces. Most inorganic and organic fluxes can tolerate the high temperatures required to supply heat to larger members. The rosin-base fluxes can quickly degrade, as indicated by the formation of thick, blackened residues on the surface. Heat and solder are removed when an adequate fillet forms at the joint opening, in order to ensure complete filling of the joint gap. The hot dip coating of parts can be performed by ultrasonic activation, without using a flux. Ultrasonic energy is coupled to the sample through the solder bath (Fig. 22). Optimum coupling is a function of the sample geometry, power level, and workpiece-horn separation. Oxide removal does not occur through simple line-of-sight erosion by the horn (Ref 28). Rather, the oxide is disrupted by the ultrasonic energy that is transferred from the horn into the substrate. Therefore, hidden surfaces can also be coated by the solder. Hot-dipped coatings can be applied to large workpieces or to leads on small electronic devices. In the latter case, care should be taken to ensure that internal connections are not damaged by the ultrasonic energy. This technology has been used to assemble tube-socket joints on heat-exchanger equipment (Ref 29).

FIG. 22 ULTRASONIC SOLDERING BATH EQUIPMENT

An important component of the solder assembly process is the fixturing that is used to support the workpieces being assembled. Fixturing details are most critical for global heating processes in which the fixture will also be raised to the soldering temperature. The purpose of fixturing is to either anchor the two substrates to prevent movement while the solder is in the molten state or permit the controlled displacement of the substrates to establish specific joint gaps not achieved with the solder preform in the solid state. Lubricants include high-temperature petroleum products or solid lubricants (MoS2 and graphite). Preloads on the workpieces are provided by springs (coil or Belleville configurations). The joint gap can be established by spacers between the workpieces or particles within the joint itself (Fig. 23). The spacer materials should not cause undesirable reactions with the solder or jeopardize the strength capacity of the joint.

FIG. 23 TECHNIQUES TO MAINTAIN JOINT GAPS. (A) BUTT JOINT WITH SPACERS IN THE SOLDER. (B) BUTT JOINT WITH SPACER IN THE FIXTURE. (C) LAP JOINT CONSTRUCTION

Fixture materials should be carefully assessed. Although high vapor pressures generated by fixture metals is not of concern at the relatively low temperatures used in soldering, factors such as thermal expansion and fixture size should be considered. For example, a CTE mismatch between the workpiece and the fixture materials, with improperly designed dimensions, can cause the joint gap to grow or close at soldering temperatures. Fixtures with large thermal mass will lengthen the heating and cooling times for the workpiece, leading to possible flux degradation, base metal or coating erosion, and longer production time. Preferred fixturing materials include oxide ceramics (alumina, beryllia, or mullite), refractory metals (molybdenum, tungsten, or tantalum), and well-oxidized steels (low-carbon and stainless steels). These

materials will generally not be wetted by solder (and flux) spillage and can accommodate the elevated temperatures. Fixtures made of machined metals, such as aluminum, copper alloys, or steels, should be annealed in order to relieve any residual stresses that arise from the stock material or machining operations. These internal stresses may cause misalignment of the workpieces or the binding of moving parts, caused by warpage generated at soldering temperatures. Preoxidized surfaces will prevent inadvertent wetting by the solder. Fixtures should be cleaned of organic residues to limit their outgassing in vacuum furnaces, prevent deposition of the contaminants onto the solderable surfaces, and limit spatter by their volatilization at soldering temperatures. Materials that contain or are coated with cadmium or zinc should not be used, because their high vapor pressures (even at soldering temperatures) may cause them to contaminate the joint area, as well as to poison vacuum and inert-atmosphere furnace systems. The workpiece temperature should be monitored, at least in the prototype stages of process development, in order to document temperature conditions. In furnace operations, it is preferred that the controlling thermocouple contact the substrates to confirm the desired temperature profile at the joint area. The relatively low temperatures of soldering allow the use of inexpensive thermocouples, such as types J, T, and K, for many furnace cycles. The thermocouples should not interfere with the fixtures or the filling of the joint. An undesirable spread of the molten solder on the substrate surface can be prevented by a number of "solder-stop" products. High-temperature tapes, which are popular in the electronics industry to prevent solder wetting of areas on printed circuit boards, are suited for many structural applications. Another product is a slurry made from powders that are mixed with water or alcohol and painted on the surface areas that are to be free of solder. The vehicle evaporates, leaving a coating that prevents spreading by the solder. Detailed descriptions of the individual soldering processes, equipment, and applications are provided in the Section "Solid-State Welding, Brazing, and Soldering Processes" in this Volume.


28. P. VIANCO AND F. HOSKING, ANALYSIS OF ULTRASONIC TINNING, NEPCON WEST, FEB 1992, P 1718 29. J. SCHUSTER AND R. CHILKO, ULTRASONIC SOLDERING OF ALUMINUM HEAT EXCHANGERS, WELD. J., OCT 1975, P 711 General Soldering Paul T. Vianco, Sandia National Laboratories

Postassembly Cleaning Procedures After the soldering operation, the workpiece is cleaned, primarily to remove flux residues that can cause corrosion of the part while in storage or during service. Other cleaning procedures include the removal of solder-stop materials, as well as stray solder particles, which can interfere with mechanical or electrical performance of the assembly. Flux residues should be removed as soon as possible after the soldering process, because their ability to be removed decreases with time, whereas their tenacity and potential for corrosive damage increase with time. Cleaning fluids and solvents to be used are determined by the particular flux residues. Guidelines are presented in the section "Fluxes" in this article. The selection of organic solvents is rapidly changing as chlorofluorocarbon materials are phased out by environmental regulations and new materials (and processes) become qualified as replacements. Some general practices must be considered when establishing cleaning procedures. First, assess the compatibility of substrate materials and filler metals with the cleaning solutions. Organic solvents are benign toward metal surfaces. Alkaline solutions used to neutralize strongly acidic fluxes can affect some base-metal finishes (for example, copper

alloys, iron alloys, and some steels). When in doubt, samples of the substrate should be exposed to the cleaning agent prior to use on final assemblies. Second, determine whether the cleaning solutions leave undesirable residues. Mineral deposits from tap water can corrode or stain substrate surfaces (especially in the continued presence of water vapor). Tap-water rinses should be followed with rinses in either deionized or distilled water. All traces of water can be removed by a final alcohol rinse. Isopropyl alcohol is generally used. Denatured alcohol and acetone should be avoided, because they also can leave residues. Third, the drying of solvent or other cleaning-solution residues should utilize dry, clean gas, such as bottled or cryogenic nitrogen gas. Compressed "house" air may contain water particles or compressor oils that can quickly recontaminate the workpiece. Fourth, limit the contact of the assembly with oily rags and fingerprints. Many instances of cosmetic staining or pitting of the workpiece surface have been traced to fingerprints. Fifth, test for the effectiveness of postcleaning operations. Unlike the well-specified procedures used by the electronics industry, such procedures are not well standardized for structural applications. Temperature-humidity chambers can be used to assess the propensity for corrosion on the workpiece (a destructive test). Cleaning effectiveness can be enhanced by thermal and mechanical assistance. Cleaning solvents and solutions have higher solubility for residues at elevated temperatures. Caution should be observed when heating solutions because of the generation of vapors that can result in health or fire hazards. The use of solvent vapors at their boiling point in vapor degreasers can remove residues from remote locations on the workpiece. However, the solvents that have been popular for vapor degreasing are being restricted from use by environmental statutes. Mechanical agitation of cleaning solutions is obtained by ultrasonic activation, high-energy sprays, and manual scouring procedures. Ultrasonics are very effective for loosening residues, particularly in hidden locations. Although generally safe for the cleaning of structural members, care must be exercised when using ultrasonics on electronic assemblies because of the possible damage to internal connections. The use of sprays or jets to force the cleaning solution into crevices and hidden areas of the workpiece can increase cleaning efficiency. Batch and in-line equipment based on spray and jet technology is currently available. Because the cleaning material passes through a jet, aerosols and mists are generated, which may create an explosion hazard. Manual scouring can remove residues on exposed surfaces only. Cleaning with sandpaper or vapor blasting metal surfaces with abrasive particles should be avoided in the postprocess cleaning steps for three reasons: First, the base-metal oxide layer protects the surface from corrosion or excessive oxidation later in service. For example, stainless steels are particularly susceptible to corrosion attack after abrasive treatments, particularly those that use steel wool or a steel brush. Second, solders are generally much softer than the base metals. Therefore, inadvertent damage can be easily done to the joint fillets, possibly jeopardizing monotonic strength and fatigue resistance. Third, abrasive grit particles can become embedded in the substrate and, particularly, in the softer solder. Dislodged particles can damage mechanical actuators that are part of the soldered assembly. The abrasive particles can also deteriorate surface solderability of the workpiece during subsequent assembly, repair, or rework procedures. Damage occurs readily to circuit boards (solder masks, coatings, and the laminate itself) by abrasive particles. Finally, storage of the parts must be considered. The extent of storage control depends on factors such as the type of assembly and its service requirements, the cost of rework, repair or scrapping of damaged parts, and the environment of the factory. Acute contamination or corrosion of finished parts can be prevented by their enclosure in bags. Popular containers are polyethylene plastic bags, which can provide short-term storage ( PA + PH + PB

(EQ 2)

where Pa is the atmosphere pressure, Ph is the hydrostatic pressure, and Pb is the pressure increase that is due to the curvature of the pore. For underwater welding, Ph is the controlling term, because it is directly related to depth. The effect of water pressure on the formation of porosity in underwater gravity welding with electrodes of the ilmenite type, high titanium oxide type, and iron powder/iron oxide type and with SM41 steel base metals has been reported (Ref 22). The amount of porosity increased with water pressure (depth) (Fig. 7). The pore shape changed from spherical to long and narrow with increasing pressure, which indicates a change of the pore-formation mechanism. In wet underwater welds, the pores contained approximately 96 vol% H2, 0.4 vol% CO, and 0.06 vol% CO2. Other bubble compositions (Ref 6, 23) of 62 to 82% H2, 11 to 24% CO, and 4 to 6% CO2 have been reported (Ref 24). More-specific data on a rutile iron powder (E7014) electrode are 45% H2, 43% CO, 8% CO2, and 4% other. The variation in bubble composition indicates the broad nature of the chemical processes associated with wet underwater welding and is most likely due to variations in electrode covering compositions, energy input, and water depth.

FIG. 7 EFFECT OF WATER PRESSURE ON POROSITY. SOURCE: REF 22

Two possible methods of reducing weld-metal porosity with depth, in the case of wet welding, are to add to the weld pool specific elements that getter the hydrogen by forming hydrides and to alter the welding parameters to reduce pore formation. At increasing travel speeds, the number of pores per volume of weld deposit goes through a maximum. A relationship between the product of travel speed (S) and bead area (A) and the amount of gas absorbed per unit volume of weld metal exists (Ref 25). This relationship suggests that for decreasing values of the product (SA), the gas absorbed per unit volume of weld metal increases for the same gas partial pressure. With small beads and a slow welding speed, the amount of gas absorbed per unit volume of weld metal is high. The faster freezing rate of the smaller, short-circuiting arc weld allows less time for gas desorption and causes more gas bubbles to be trapped before floating to the surface of the weld pool. This relationship holds significance for many of the practices used in wet underwater welding, which produce small weld deposits. In underwater welding, the relationship between porosity and welding current is strongly influenced by the type of moisture resistance coating on the electrode. The amount of gas pickup increases as the arc was lengthened (Ref 21). In general, the hydrogen absorption and, therefore, porosity levels in welding can be minimized by using a low current with direct current electrode positive (DCEP), a high current with direct current electrode negative (DCEN), a short arc, and a fast travel speed.


6. K. MASUBUCHI, UNDERWATER FACTORS AFFECTING WELDING METALLURGY, PROC. CONF. UNDERWATER WELDING OF OFFSHORE PLATFORMS AND PIPELINES, AWS, 1980, P 8198 21. R.E. TREVISAN, D.D. SCHWEMMER, AND D.L. OLSON, THE FUNDAMENTALS OF WELD METAL PORE FORMATION, WELDING: THEORY AND PRACTICE, D.L. OLSON, R.D. DIXON, AND A.L. LIBY, ED., ELSEVIER SCIENCE B.V., 1990 22. Y. SUGA AND H. ATSUSHI, "ON THE FORMATION OF POROSITY IN UNDERWATER WELD METAL (THE 1ST REPORT)--EFFECT OF WATER PRESSURE ON FORMATION OF POROSITY," IIW DOC. IX-1388-86, AMERICAN COUNCIL, AWS, 1986 23. N.M. MALATOR, THE PROPERTIES OF THE BUBBLES OF STEAM AND GAS AROUND THE ARC IN UNDERWATER WELDING, AUTOMAT. WELD., VOL 18 (NO. 12), 1965, P 25-29 24. E.A. SILVA, GAS PRODUCTION AND TURBIDITY DURING UNDERWATER SHIELDED METALARC WELDING WITH IRON POWDER ELECTRODES, NAV. ENG. J., VOL 12, 1971 25. G.R. SALTER AND D.R. MILNER, GAS METAL REACTIONS IN ARC WELDING, BRIT. WELD. J., VOL 33, 1965

Underwater Welding S. Liu and D.L. Olson, Colorado School of Mines, S. Ibarra, Amoco Corporation Research

Fatigue as a Function of Porosity The designers of offshore structures are usually concerned with the initiation and growth of fatigue cracks in fracturecritical members, because of noted decreases in fracture toughness, as well as the presence of varying amounts of porosity. Cyclic fatigue stresses will develop with any natural sea movement and will increase during storms. Drilling the tip of a fatigue crack is a common field technique that is used to prevent further crack growth. In the same manner, the presence of porosity in weldments can be beneficial, because the pores act as pinning sites for fatigue cracks and can either retard or stop their growth. In particular, the presence of some amount of porosity retarded fatigue crack growth at low stress in surface habitat and wet underwater welds (Ref 26) (Fig. 8). The low-porosity dry habitat weld was shown to contain approximately 3% porosity on the fatigue fracture surface, whereas the high-porosity wet underwater weld contained approximately 12% porosity on the fatigue fracture surface.

FIG. 8 COMPARISON OF EFFECTS OF POROSITY AND FREQUENCY ON THE FATIGUE CRACK GROWTH BEHAVIOR OF WELDS TESTED IN SEAWATER. SOURCE: REF 26

In comparison to normal wrought ferrite-pearlite steels, all experimental underwater welds exhibited lower growth rates for low values of ∆K, where ∆K is the cyclic variation in the stress intensity factor, K. However, the wet welds tended to have a higher fatigue crack growth rate than the wrought alloys at high values of ∆K, particularly in wet welds. At low ∆K values, the wet underwater welds have the most fatigue crack growth resistance. The large number of pores act as pinning sites for the advancing fatigue crack front. Increased porosity in a weldment yielded lower fatigue crack growth at low ∆K values.

At high ∆K values, the crack growth rate is much greater for the higher-porosity content. At higher stress levels, the mechanical behavior of the sample is more like that experienced in a tensile test. The high numbers of pores act to reduce the cross-sectional area and, thus, allow the cracks to propagate at a high rate. The use of wet weld repairs in areas of high cyclic stress should be accompanied by caution and a load-reducing joint design.


26. D.K. MATLOCK, G.R. EDWARDS, D.L. OLSON, AND J. IBARRA, AN EVALUATION OF THE FATIGUE BEHAVIOR IN SURFACE, HABITAT, AND UNDERWATER WET WELDS, UNDERWATER WELDING, PERGAMON PRESS, 1983, P 303-310 Underwater Welding S. Liu and D.L. Olson, Colorado School of Mines, S. Ibarra, Amoco Corporation Research

Heat Sources With increasing depth and, thus, pressure, the welding parameter space for sustaining a welding arc becomes more restricted. This effect necessitates a more careful selection of the welding voltage (gap) and current for a given size and type of welding electrode. The decreasing welding parameter space also reflects a shift of current to lower values and voltage to higher values (see Fig. 3), which changes the nature of the heat input. Electrode coating thickness can also influence the arc characteristics (Ref 16). Figure 9 illustrates that increasing coating thickness also raises the arc characteristic curves.

FIG. 9 EFFECTS OF DEPTH ON ARC VOLTAGE FOR THREE DIFFERENT GTAW ELECTRODE DIAMETERS. SOURCE: REF 29

With an increase in the welding depth, the duration and frequency of short circuits and arc extinctions were reported to grow significantly (Ref 28), and the process stability deteriorated. The observed rise of welding current is associated with the increase in the number of short circuits. The addition of rare earth metals into the charge of rutile-base flux-cored wire

provides considerable improvement in arc stability and reduces arc downtime (arc extinctions and short circuits) by nearly two times. The more-stable arc burning that occurs when welding with the rare earth metal-containing wire allows its use in any spatial position and within the depth ranges investigated. An approach that separates the wet conditions from the weld pool has been demonstrated, in which a welding torch that produces a local gas cavity by controlling the fluid flow in the environment of the weld is utilized (Ref 29). This method hinders the water invasion by the momentum of the fluid jet and by the smooth exhausting of cavity gases by flow entrapment. The process allows near-surface welding conditions to exist and permits the flexibility of wet welding practices while achieving weld quality that is similar to dry-environment underwater welding.


16. V.Y. KONONENKO, EFFECT OF WATER SALINITY AND MECHANIZED UNDERWATER WELDING PARAMETERS ON HYDROGEN AND OXYGEN CONTENT OF WELD METAL, WELDING UNDER EXTREME CONDITIONS, PERGAMON PRESS, 1989, P 113-118 28. I.K. POKHODNYA, V.N. GORPENYUK, V.YA. KONONENKO, V.E. PONOMAREV, AND S.YU. MAKSIMOV, SOME PECULIARITIES OF ARC BURNING AND METAL TRANSFER IN WET UNDERWATER SELF-SHIELDING FLUX-CORED WIRE WELDING, WELDING UNDER EXTREME CONDITIONS, PERGAMON PRESS, 1989, P 151 29. S. FUKUSHIMA, T. FUKUSHIMA, AND J. KINUGAWA, UNDERWATER WET PLASMA WELDING IN PRESSURIZED WATER, TRANS. MET. INST. MET., VOL 19 (NO. 3), 1977, P 133-151 Underwater Welding S. Liu and D.L. Olson, Colorado School of Mines, S. Ibarra, Amoco Corporation Research

Practical Applications for Underwater Welding Underwater welding techniques are used primarily for the repair of offshore platforms, particularly after the normal 5 year inspection period. However, these techniques have also been very useful during the installation of new offshore structures and undersea pipelines, the installation of hot taps, the repair of dock and harbor facilities, the modification of and addition to underwater structures, and the repair of nuclear facilities. Examples of maintenance and repair applications are described below. Offshore Structures. Underwater welding techniques are used to repair damages caused by corrosion or by fatigue

cracking, and to repair members that have been damaged by ship impact. The repair or replacement of structural members damaged during installation by objects that have fallen overboard or by other accidents have also been reported. The underwater wet welding repair of these structures requires that the carbon equivalent of the steel be below 0.40 wt% (as calculated using the International Institute of Welding formula) to prevent hydrogen cracking. For steels with high carbon equivalents, the use of multiple temper bead techniques is required in order to reduce the hardness in the HAZ and eliminate the presence of hydrogen. Undersea Pipelines. The repair or replacement of undersea pipelines is usually conducted in a hyperbaric dry chamber.

Although sleeve repairs using wet welds have been tested successfully, their use has not been reported to date. Hyperbaric chambers are used regularly for underwater tie-ins and for repairs of defects. Harbor Facilities. Corrosion and collision damage to sheet and pilings have been repaired. The repair and replacement

of tubular dock braces, dock supports, and tanker moorings have also been reported. Floating Vessels. Permanent and temporary repairs to holes in ship and barge hulls have been conducted. The hulls and

pontoons of semisubmersible drill ships have also been repaired.

Nuclear Power Plants. The wet welding of nuclear power plant components is usually conducted with austenitic

stainless steel electrodes in order to match the austenitic base metal used in the construction of these power plants. Damaged reactor internal components and leaks in pool liners are the usual types of repairs. Specifications. The first underwater welding specification, known as ANSI/AWS D3.6, "Specification for Underwater

Welding," was published in 1983 by the American Welding Society. This specification provides comprehensive information that enables engineers to select the welding method (wet or dry) that meets the established fitness-for-purpose requirements. The specification also provides a means for the engineer to make critical decisions in structural fabrication when underwater welding techniques are used. It is also a performance specification that is intended to assist the purchaser in specifying and obtaining the desired mechanical properties when using underwater welding techniques. The specification also allows contractors to prepare estimates and quotations for each project with a clear understanding of what is required. The original specification has been revised and reissued several times. The AWS D3.6 specification defines four weld types, or classes. Each weld type has a specific set of quality requirements that must be verified during procedure qualification. The type of weld selected for a specific application is determined by the user. Type A welds are represented by a set of requirements that produces an underwater weld comparable to a surface weld.

Underwater welds that meet these requirements are usually hyperbaric (dry) welds. Type B welds are usually reserved for wet underwater welds. They are intended for less-critical applications where

reduced ductility and increased porosity can be tolerated. Typical applications include scallop sleeves on offshore structures. Type C welds satisfy lesser requirements than type B welds and are intended for applications where load-beating

function is not a primary consideration. Type O welds usually meet the requirements of surface dry welds, as well as those of other codes and specifications. Other underwater welding documents are developed by agencies such as the International Institute of Welding (IIW

Doc SCUW 124-90), the American Society of Mechanical Engineers (ASME Section X/IWA 4550 for Underwater Welding of Nuclear Vessels), and the Norwegian agency Det Norske Veritas (DNV) (Document RP3604, Recommended Practices for Underwater Welding). Underwater Welding S. Liu and D.L. Olson, Colorado School of Mines, S. Ibarra, Amoco Corporation Research

References

1. C.L. TSAI AND K. MASUBUCHI, INTERPRETIVE REPORT ON UNDERWATER WELDING, WELD. RES. COUNC. BULL., VOL 224, 1977, P 1-37 2. N. CHRISTENSEN, THE METALLURGY OF UNDERWATER WELDING, UNDERWATER WELDING, PERGAMON PRESS, 1983, P 71-79 3. A. HAUSI AND Y. SUGA, ON COOLING OF UNDERWATER WELDS, TRANS. JPN. WELD. SOC., VOL 2 (NO. 1), 1980 4. C.L. TSAI AND K. MASUBUCHI, MECHANISMS OF RAPID COOLING IN UNDERWATER WELDING, APPL. OCEAN RES., VOL 1 (NO. 2), 1979, P 99-110 5. C.E. GRUBBS AND O.W. SETH, "MULTIPASS ALL POSITION WET WELDING, A NEW UNDERWATER TOOL," OFFSHORE TECHNOLOGY CONFERENCE, OTC 1620 (HOUSTON), 1972 6. K. MASUBUCHI, UNDERWATER FACTORS AFFECTING WELDING METALLURGY, PROC. CONF. UNDERWATER WELDING OF OFFSHORE PLATFORMS AND PIPELINES, AWS, 1980, P 8198

7. N.M. MADATOR, INFLUENCE OF THE PARAMETERS OF THE UNDERWATER WELDING PROCESS ON THE INTENSITY OF METALLURGICAL REACTIONS, WELD. RES. ABROAD, VOL 3, 1972, P 63 8. O. GRONG, D.L. OLSON, AND N. CHRISTENSEN, ON THE CARBON OXIDATION IN HYPERBARIC MMA WELDING, MET. CONSTR., VOL 17 (NO. 12), 1985, P 810R-814R 9. D.L. OLSON AND S. IBARRA, UNDERWATER WELDING METALLURGY, PROC. FIRST OMAE SPECIALTY SYMPOSIUM ON OFFSHORE AND ARCTIC FRONTIERS, ASME, 1986, P 439-447 10. S. IBARRA, C.E. GRUBBS, AND D.L. OLSON, FUNDAMENTAL APPROACHES TO UNDERWATER WELDING METALLURGY, J. MET., VOL 40 (NO.12), 1988, P 8-10 11. S. IBARRA AND D.L. OLSON, WET UNDERWATER STEEL WELDING, FERROUS ALLOY WELDMENTS, VOL 69-70, KEY ENG. MATERIALS, TRANS TECH PUBLICATIONS, ZURICH, 1992, P 329-378 12. O. OZAKI, J. NAIMAN, AND K. MASUBUCHI, A STUDY OF HYDROGEN CRACKING IN UNDERWATER STEEL WELDS, WELD. J., VOL 56 (NO. 8), 1977, P 231S-237S 13. W.A. STALKER, P.H.M. HART, AND G.R. SALTER, "AN ASSESSMENT OF SHIELDED METAL ARC ELECTRODES FOR THE UNDERWATER WELDING OF CARBON MANGANESE STRUCTURAL STEELS," OFFSHORE TECHNOLOGY CONFERENCE, OTC REPORT 2301 (HOUSTON), 1975 14. B. CHEW, PREDICTION OF WELD METAL HYDROGEN LEVELS OBTAINED UNDER TEST CONDITIONS, WELD. J., VOL 52, 1973, P 386S-391S 15. L.I. SOROKIN AND Z.A. SIDLIN, THE EFFECT OF ALLOYING ELEMENTS AND OF MARBLE IN AN ELECTRODE COATING ON THE SUSCEPTIBILITY OF A DEPOSITED NICKEL CHROME METAL TO PORE FORMATION, SVAR. PROIZVOD., NO. 11, 1974, P 7-9 16. V.Y. KONONENKO, EFFECT OF WATER SALINITY AND MECHANIZED UNDERWATER WELDING PARAMETERS ON HYDROGEN AND OXYGEN CONTENT OF WELD METAL, WELDING UNDER EXTREME CONDITIONS, PERGAMON PRESS, 1989, P 113-118 17. H. HOFFMEISTER AND K. KUSTER, PROCESS VARIABLES AND PROPERTIES OF UNDERWATER WET SHIELDED METAL ARC LABORATORY WELDS, UNDERWATER WELDING, PERGAMON PRESS, 1983, P 115-125 18. K. HOFFMEISTER AND K. KUSTER, PROCESS VARIABLES AND PROPERTIES OF WET UNDERWATER GAS METAL ARC LABORATORY AND SEA WELDS OF MEDIUM STRENGTH STEELS, UNDERWATER WELDING, PERGAMON PRESS, 1983, P 239-246 19. A. MATSUNAWA, H. TAKEMATA, AND I. OKAMOTO, ANALYSIS OF TEMPERATURE FIELD WITH EQUIRADIAL COOLING BOUNDARY AROUND MOVING HEAT SOURCES--HEAT CONDUCTION ANALYSIS FOR ESTIMATION OF THERMAL HYSTERESIS DURING UNDERWATER WELDING, TRANS. JWRI, VOL 9 (NO. 1), 1980, P 11-18 20. K. OLSEN, D.L. OLSON, AND N. CHRISTENSEN, WELD BEAD TEMPERING OF HEAT AFFECTED ZONE, SCAND. J. METALL., VOL 11, 1982, P 163-168 21. R.E. TREVISAN, D.D. SCHWEMMER, AND D.L. OLSON, THE FUNDAMENTALS OF WELD METAL PORE FORMATION, WELDING: THEORY AND PRACTICE, D.L. OLSON, R.D. DIXON, AND A.L. LIBY, ED., ELSEVIER SCIENCE B.V., 1990 22. Y. SUGA AND H. ATSUSHI, "ON THE FORMATION OF POROSITY IN UNDERWATER WELD METAL (THE 1ST REPORT)--EFFECT OF WATER PRESSURE ON FORMATION OF POROSITY," IIW DOC. IX-1388-86, AMERICAN COUNCIL, AWS, 1986 23. N.M. MALATOR, THE PROPERTIES OF THE BUBBLES OF STEAM AND GAS AROUND THE ARC IN UNDERWATER WELDING, AUTOMAT. WELD., VOL 18 (NO. 12), 1965, P 25-29 24. E.A. SILVA, GAS PRODUCTION AND TURBIDITY DURING UNDERWATER SHIELDED METALARC WELDING WITH IRON POWDER ELECTRODES, NAV. ENG. J., VOL 12, 1971 25. G.R. SALTER AND D.R. MILNER, GAS METAL REACTIONS IN ARC WELDING, BRIT. WELD. J., VOL 33, 1965

26. D.K. MATLOCK, G.R. EDWARDS, D.L. OLSON, AND J. IBARRA, AN EVALUATION OF THE FATIGUE BEHAVIOR IN SURFACE, HABITAT, AND UNDERWATER WET WELDS, UNDERWATER WELDING, PERGAMON PRESS, 1983, P 303-310 27. G. EDWARDS, C.J. ALLUM, B.E. PINFOLD, AND J.H. NIXON, THE EFFECT OF PRESSURE ON THE TUNGSTEN ARGON WELDING ARC, PROC. CONF. PHYSICS AND WELD POOL BEHAVIOR, WELDING INSTITUTE, 1979, P 101-107 28. I.K. POKHODNYA, V.N. GORPENYUK, V.YA. KONONENKO, V.E. PONOMAREV, AND S.YU. MAKSIMOV, SOME PECULIARITIES OF ARC BURNING AND METAL TRANSFER IN WET UNDERWATER SELF-SHIELDING FLUX-CORED WIRE WELDING, WELDING UNDER EXTREME CONDITIONS, PERGAMON PRESS, 1989, P 151 29. S. FUKUSHIMA, T. FUKUSHIMA, AND J. KINUGAWA, UNDERWATER WET PLASMA WELDING IN PRESSURIZED WATER, TRANS. MET. INST. MET., VOL 19 (NO. 3), 1977, P 133-151 Welding for Cryogenic Service T.A. Siewert and C.N. McCowan, National Institute of Standards and Technology

Introduction CRYOGENIC TEMPERATURES cause many structural alloys to become brittle, which is an unacceptable condition in most structural applications. Therefore, structures built for service at low temperatures are typically made from alloys that maintain some ductility at the service temperatures. Cryogenic alloys include 9Ni steels, austenitic stainless steels, manganese stainless steels, maraging steels, titanium, aluminum, and nickel alloys. The choice of weld-metal alloy may depend solely on the strength of the alloy at a given temperature or on a combination of strength, toughness, fatigue resistance, thermal conductivity, magnetic permeability, and other considerations. Wrought or cast alloy structural members are usually joined using filler materials with compositions that are similar, but not identical, to the parent material. The weld composition is optimized to compensate for the inherent differences in the properties of weld and parent material that are due to grain structure, inclusion content, and cooling rate differences. The goal is to match, as closely as possible, the mechanical and physical properties of the weld and parent material. When nonmatching weld compositions are most appropriate, differences between the welds and parent material in terms of thermal contraction, corrosion, and other factors must be considered. This article discusses these differences and explains how they affect the choice of the weld filler metal. Welding for Cryogenic Service T.A. Siewert and C.N. McCowan, National Institute of Standards and Technology

Effects of Cryogenic Service on Properties Strength. As the temperature drops below room temperature, some cryogenic alloys strengthen appreciably. Unless

some other property degrades and offsets this effect, such strengthening is very beneficial in reducing the required weight and thickness of structural members that are designed on the basis of load. Figure 1 shows this strengthening effect for several alloy systems, using data that are representative of both base and weld metal. Unfortunately, not all alloys strengthen at the same rate. Therefore, a comparison of the strengths at one temperature cannot be used to predict the relative strengths at a different temperature. Most literature data are for certain convenient test temperatures (4 K, 77 K, and 298 K). Data for intermediate temperatures indicate that strength is a smoothly varying function of temperature, whereas data for the three common test temperatures can be linked with curved lines. Because the behavior is very predictable, this discussion concentrates on the extreme temperature, 4 K, where the effects are most pronounced. Behavior at 77 K can be estimated by interpolation.

FIG. 1 YIELD STRENGTH VERSUS TEMPERATURE (4 TO 300 K) FOR VARIOUS BASE METALS AND WELD METALS THAT HAVE BEEN STUDIED FOR CRYOGENIC APPLICATIONS

As shown in Fig. 1, some alloys can reach very high yield strengths (well above 1000 MPa, or 140 ksi) at 4 K. Others, such as aluminum alloys, start from moderate or low strengths and respond mildly to a lowering of temperature (Ref 1, 2). Nine-percent nickel (Fe-9Ni) starts with a strength that is high at room temperature and rises from that point (Ref 1, 2). High-manganese steels, such as Fe-22Mn, and 21Cr-6Ni-9Mn stainless steel at 4 K approach the ultra-high-strength levels of 9Ni steel, maraging steel, and titanium (Ref 3, 4). The austenitic (300 series) stainless steels offer medium strength at room temperature, but strengthen with decreasing temperature at a rate dependent on the nitrogen content of the alloy. As shown in Fig. 1, the strengthening of these alloys is a strong function of their nitrogen content: the trend line for the alloys with a higher nitrogen content shows much greater strengthening than the trend line for the alloys with a lower nitrogen content (Ref 5). The maximum strength that can be reached is a function of the solubility limit of nitrogen in the stainless steel. This limit, in turn, is a function of composition and other factors. To control strength values, the nitrogen content of the weld must be carefully controlled, or else two effects can occur: pores may form in the weld during solidification (as the supersaturated liquid weld metal rejects nitrogen) or the weld pool may lose nitrogen through evaporation, weakening the weld. It should be noted that the strength of 21Cr-6Ni-9Mn is much higher than that of other austenitic steels because its nitrogen content is near 0.3 wt%. This higher level of strength is made possible by the improved nitrogen solubility in an alloy with a higher manganese content. Equation 1 was developed to predict the strength of welds in austenitic stainless steel at a temperature of 4 K:

σY(MPA) = 180 + 3200 N + 33 Mo + 32 Mn + 13 Ni

(EQ 1)

where the elemental symbols represent their concentration in wt% (Ref 5). Equation 1 shows which elements have the greatest effect on strength and is very similar, in terms of elements and their coefficients, to an equation developed to predict the strength of austenitic stainless steel wrought material (Ref 6). Thus, it can be concluded that:

• • •

AUSTENITIC WELD STRENGTH CAN BE ADJUSTED BY CHANGING THE COMPOSITION TO MEET THE NEEDS OF THE APPLICATION. WELD STRENGTH IS DETERMINED BY THE SAME FACTORS AS THOSE THAT DETERMINE THE BASE METAL STRENGTH. MATCHING THE COMPOSITION OF THE WELD TO THAT OF THE WROUGHT MATERIAL WILL RESULT IN NEARLY EQUIVALENT STRENGTHS.

Fracture Toughness at 4 K. Because many cryogenic structures are subject to fracture and other failure mechanisms,

most designs must consider toughness, as well as strength. Unfortunately, toughness often degrades as the strength is increased or as the temperature decreases. Figure 2 shows data for strength and fracture toughness at 4 K for some alloys that have been either proposed or actually used in cryogenic applications. The plotting of yield strength versus fracture toughness, KIc, provides an easy way to evaluate alloys, because these mechanical properties are common design criteria. A trend band or border was drawn around the data when warranted by the quantity of data. For some alloys, the data were so sparse that individual data points were plotted. Although most of these data are for base-metal alloys, data for selected welds are included to show that base-metal properties for important cryogenic alloys can be met or exceeded for high-toughness, fracture-critical applications.

FIG. 2 YIELD STRENGTH VERSUS FRACTURE TOUGHNESS FOR VARIOUS BASE METAL ALLOYS AND SELECTED WELDS AT 4 K

Data for a Ti-6Al-4V alloy and an electron-beam (EB) weld showed an abrupt decrease in toughness at test temperatures below 125 K; toughness was lowest at the boundary between the heat-affected zone (HAZ) and the fusion zone of the weld (Ref 7). These data are grouped in the lower right corner of Fig. 2. Generally, titanium alloys and other ultrahighstrength alloys, such as 18Ni maraging steel or 9Ni steel, exhibit low toughness at 4 K (as might be expected of an alloy with very high strength). Therefore, these alloys are most commonly used for compression loading or other applications where their limited ductility is acceptable (Ref 8, 9, 10). The 18Ni maraging steel and the titanium alloy do not show a transition from ductile to brittle fracture at low temperatures, as do 9Ni and other lower-nickel-content steels (Ref 1, 2, 9). No data for weld toughness were found for 18Ni maraging steels at 4 K. Data for aluminum alloys and some welds are also included in Fig. 2 (Ref 11, 12). These data represent traditional aluminum alloys, as well as some of the new aluminum-lithium alloys. The zones indicated in Fig. 2 for the aluminum

and titanium alloys are intended to loosely identify mechanical property regions for the alloys, rather than to define limits for design purposes. At a temperature of 4 K, considerable data exist for type 316L stainless steel wrought alloy (a common cryogenic structural material). Therefore, Fig. 2 includes a trend band for yield strength versus toughness (Ref 6). This alloy has a good balance of strength and toughness, and has found application as the structural case for large superconducting magnets. As its strength increases (through additions of nitrogen), its toughness decreases. This variation in strength with nitrogen content was shown in Fig. 1 as curves for two different nitrogen contents, whereas Fig. 2 shows the effect in terms of toughness as a function of strength at constant temperature. This trend shows the interrelationships between properties that must be considered by structural designers. Although the trend line for the matching-composition electrodes is not included in Fig. 2, it extends over the same strength range, but at only 50 to 70% of the toughness of the base metal (Ref 4). This lower toughness for welding electrodes of matching composition has driven the search for electrodes of different compositions that can match the properties of these base metals. Two fully austenitic stainless steel weld compositions (shown as A and B in Fig. 2 and described in more detail in Table 1) have been proposed for joining AISI 316LN stainless steel (Ref 13, 14, 15).

TABLE 1 COMPOSITIONS FOR THE FOUR WELDS (A THROUGH D) INCLUDED IN FIG. 2

WELD COMPOSITION, % FE CR NI MN A BAL 13 5 22 B BAL 20 25 1.6 C BAL 18.1 20.4 5.4 D 1 20 BAL 3

PROCESS N 0.21 0.16 0.16 ...

C 0.04 0.01 0.03 0.02

S 0.004 0.001 0.007 ...

P 0.013 0.010 0.006 ...

EBW GMAW GMAW GTAW

Two other iron-base alloys, Fe-22Mn and 21Cr-6Ni-9Mn, have been used in structures that require high strength and toughness at 4 K. Figure 2 includes a trend line for a high-manganese alloy, Fe-22Mn, which is one of several ironmanganese alloys that have been developed since the early 1980s (Ref 3, 16). Its particular advantage is an ability to attain a strength higher than that of type 316 stainless steel. A matching-composition weld, shown as C in Fig. 2, was produced by electron-beam welding (EBW) (Ref 17). The strength of the weld was approximately 5% lower than that of the parent plate material, presumably because of the loss of nitrogen in the welding process. However, the fracture toughness of the weld exceeded the toughness of the plate material. Matching-composition welds have also been made using shielded metal arc welding (SMAW), gas-metal arc welding (GMAW), gas-tungsten arc welding (GTAW), and submerged arc welding (SAW) processes. The toughness of the welds produced by these other processes is lower than that of the EB weld, because their inclusion contents are higher (Ref 16). Another common cryogenic alloy is 21Cr-6Ni-9Mn, for which one value is included in Fig. 2 (Ref 2). A trend band has not been developed, but the data would be expected to have a similar trend to those of Fe-22Mn and 316LN (but extending to a higher strength than that of 316LN, because the 21Cr-6Ni-9Mn can contain up to 0.4 wt% N). A nickelbase weld composition (weld D in Fig. 2) has been proposed for 21Cr-6Ni-9Mn. Although this nickel-base weld undermatches the strength, it overmatches the toughness. Thus, welds can be made in locations of lower stress, which permits the use of this high-strength alloy in the construction of structures. Table 1 lists the compositions of the four welds included in Fig. 2. It shows that these welds (other than the Fe-22Mn composition) have more nickel than the iron-base alloys that they join. These weld composition data represent the best combinations of strength and toughness reported for the alloys. Because the matching-composition welds tend to have lower toughness than the parent plate, these nonmatching compositions are being considered for fracture-critical welds. Subsequent sections in this article describe how thermal expansion and other characteristics of the nonmatching compositions might affect the suitability of these alloys in particular applications. Another common cryogenic temperature is 77 K, the normal boiling temperature of liquid nitrogen. Because toughness is higher at 77 K than at 4 K, higher-strength alloys can be considered for fracture-critical applications. Like other iron-base alloys that show a significant improvement in toughness with increasing temperature, Fe-21Cr-6Ni-9Mn is more suited for use at 77 K, where its high strength is an attractive option, when compared with that of 316LN alloy. For example, Fe-

21Cr-6Ni-9Mn base metal drops in strength from a level near 1450 MPa (210 ksi) at 4 K to near 1150 MPa (165 ksi) at 77 K, which is still a very respectable level of strength. Charpy V-Notch Impact Energy at 77 K. Although not as useful in fitness-for-service calculations as KIc fracture

toughness, the Charpy V-notch (CVN) test is a low-cost way to measure the resistance of an alloy to fracture. A literature search found that most data on fracture resistance at 77 K were reported as CVN data. In Fig. 3, CVN absorbed-energy data at 77 K are plotted versus the yield strength for a number of weld metals. The trends reiterate the effects of strength and nickel content on CVN toughness. Because many structural designs specify a minimum of 30 or 40 J (22 to 30 ft · lbf), numerous weld compositions can be considered at this temperature, including compositions that match that of the base plate.

FIG. 3 CHARPY V-NOTCH ABSORBED ENERGY VERSUS STRENGTH FOR VARIOUS WELD METALS AT 77 K

In the lower right quadrant of Fig. 3 is a data point for a matching-composition weld for the Fe-21Cr-6Ni-9Mn steel (Ref 18). The low absorbedenergy indicates that the nickel-base undermatching weld compositions should be considered for applications with this composition at 77 K (Ref 19). Nickel-base weld compositions are also used for welding the Fe-9Ni alloy included in Fig. 3 (Ref 10). For the austenitic (300 series) stainless steels, there are two trend bands. The lower band is for welds produced by the SMAW process, whereas the upper is for welds produced by the GTAW process. The difference is attributed to weld inclusions, as discussed in the next section of this article. Above the GTAW trend band are the four special weld compositions identified in Table 1. These higher absorbed energy levels are comparable to the level that type 316LN plate would develop at this temperature. Equation 2 was developed to predict the absorbed energy at 77 K for the welds in austenitic stainless steel:

CVN(J) = 19 - 1.4 FN - 890 C2 + 1.4 NI

(EQ 2)

where the elemental symbols represent the composition in wt%, and FN is a measure of the delta ferrite content (calculated by the Schaeffler diagram coefficients and allowed to be negative) in the weld (Ref 20).

This equation was not developed for weld compositions strengthened by nitrogen. Therefore, carbon represents the principal strengthening element. If carbon is considered as a generic strengthening term (combined effects of carbon, nitrogen, manganese, molybdenum, and others), then Eq 2 shows that increased strength and ferrite content lower the CVN toughness, and that nickel increases toughness. The effects of these factors on toughness are generally accepted to be true for stainless steel weld and base material. Because an increase in nickel content also decreases the ferrite content, nickel additions improve the toughness via two terms in Eq 2. This information on the effect of nickel confirms the reasons for the better toughness of the four nonmatching compositions (Ref 13, 14, 15, 17). Inclusion Content. In austenitic stainless steels and manganese stainless steels, the fracture modes of both base metals

and welds are usually quite ductile, and fracture originates at inclusions in the microstructure. A minor difference between the fracture surfaces of the base metal and weld metal is that those of the welds have a finer ductile dimple spacing. This finer spacing has been attributed to the higher inclusion content of the welds and explains why the welds are less tough than those of the matching-composition wrought material (Ref 21, 22, 23). Figure 4 shows normalized toughness data versus inclusion spacings for both wrought materials and welds (Ref 21, 24). The vertical axis is known as a quality index, and it allows alloys with different strengths and nickel contents to be compared on a scale that has fracture toughness in the numerator (higher values are better). The similarity of the trend line for both sets of data supports the need to reduce the inclusion content of welds that require a higher fracture toughness. Most of the welds in Fig. 2 were produced using welding processes and procedures that minimize the inclusion content. The exceptions here are the nickel-alloy welds and the very high nickel content (greater than 20 wt%) stainless steel welds. Preliminary data indicate that the toughness of these compositions is less sensitive to inclusion content. Therefore, these materials could be more appropriate for welding at higher deposition rates.

FIG. 4 NORMALIZED TOUGHNESS DATA (QUALITY INDEX) VERSUS INCLUSION SPACING FOR TYPE 316LN BASE METALS AND WELDS, SHOWING HOW THE WELD DATA FOLLOW THE SAME TREND AS THE BASE METAL, BUT A SMALLER INCLUSION SPACING RESULTS IN LOWER TOUGHNESS FOR THE WELDS

Fatigue Strength. Few data are available for the fatigue strength of cryogenic welds. However, the data that are available suggest that fatigue strength could limit the structural life in high-cycle applications. Table 2 lists data from a study that used the GTAW process to join a number of different materials (Ref 25). The data show that fatigue strength does not necessarily increase as temperatures decrease from 295 to 4 K. Expected yield strength data, extracted from Fig. 1, have been added for comparison.

TABLE 2 FATIGUE STRENGTH FOR GAS-TUNGSTEN ARC WELDS AT 4 × 104 CYCLES

SPECIMEN

ESTIMATED FATIGUE STRENGTH AT 295 K AT 4 K MPA KSI MPA KSI 21CR-6NI-9MN TO 304L 380 55 360 52 21CR-6NI-9MN TO 316L 420 60 390 57 21CR-6NI-9MN BASE 370 54 290 42 METAL

YIELD STRENGTH AT 4 K MPA 1050 1100 1200

KSI 150 160 175

The coefficient of thermal expansion (CTE) can have an important bearing on structural integrity. If the weld metal has a CTE that is very different from that of the base metal, cooling from room temperature to cryogenic temperature can introduce substantial stresses. These stresses can add to the residual stresses that exist around any weld and can permit the plastic deformation of the structure. Repeated thermal cycles could lead to work hardening of the microstructure or to fatigue failure, in the most severe cases. Other physical properties that are often important in cryogenic structures are phase stability and magnetic permeability.


1. MECHANICAL, THERMAL, ELECTRICAL, AND MAGNETIC PROPERTIES OF STRUCTURAL MATERIALS, HANDBOOK ON MATERIALS FOR SUPERCONDUCTING MACHINERY, 1ST ED., MCIC-HB-04, METALS AND CERAMICS INFORMATION CENTER, BATTELLE COLUMBUS LABORATORIES, JAN 1977 2. "MATERIALS FOR SUPERCONDUCTING MAGNET SYSTEMS: MECHANICAL AND PHYSICAL PROPERTIES DATA AT LOW TEMPERATURES," BELFOUR STULEN DIVISION, TRAVERSE CITY, MI, 1979 3. T. HORIUCHI, R. OGAWA, AND M. SHIMADA, CRYOGENIC FE-MN AUSTENITIC STEELS, ADV. CRYOGENIC ENG. (MATER.), VOL 32, 1985 4. R.L. TOBLER, T.A. SIEWERT, AND H.I. MCHENRY, STRENGTH-TOUGHNESS RELATIONSHIP OF AUSTENITIC STAINLESS STEEL WELDS AT 4 K, CRYOGENICS, VOL 26, 1986 5. C.N. MCCOWAN AND T.A. SIEWERT, INFLUENCE OF MOLYBDENUM ON THE STRENGTH AND TOUGHNESS OF STAINLESS STEEL WELDS FOR CRYOGENIC SERVICE, "MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES--X," NBSIR 87-3067, NATIONAL BUREAU OF STANDARDS, 1987 6. N.J. SIMON AND R.P. REED, STRENGTH AND TOUGHNESS OF AISI 304 AND 316 AT 4 K, "MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES--X," NBSIR 86-3050, NATIONAL BUREAU OF STANDARDS, 1986 7. R.L. TOBLER, LOW TEMPERATURE FRACTURE BEHAVIOR OF A TI-6AL-4V ALLOY AND ITS ELECTRON BEAM WELDS, STP 651, ASTM, 1978 8. R.L. TOBLER, R.P. REED, AND R.E. SCHRAMM, CRYOGENIC TENSILE, FATIGUE, AND FRACTURE PARAMETERS FOR A SOLUTION-ANNEALED 18 PERCENT NICKEL MARAGING STEEL, J. ENG. MATER. TECHNOL., VOL 100, APRIL 1978, P 189-194 9. J.A. WAGNER, CORRELATION OF MECHANICAL PROPERTIES WITH METALLURGICAL STRUCTURE FOR 18NI 200 GRADE MARAGING STEEL AT ROOM AND CRYOGENIC TEMPERATURES, CRYOGENICS, VOL 31, SEPT 1991 10. R.D. STOUT AND S.J. WIERSMA, FRACTURE TOUGHNESS OF MODERN 9% NICKEL CRYOGENIC STEELS, ADV. CRYOGENIC ENG. (MATER.), VOL 32, 1985 11. J. GLAZER, S.L. VERASCONI, E.N.C. DALDER, W. YU, R.A. EMIGH, R.O. RICHIE, AND J.W. MORRIS, JR., CRYOGENIC MECHANICAL PROPERTIES OF AL-CU-LI-ZR ALLOY 2090, ADV. CRYOGENIC ENG. (MATER.), VOL 32, 1985

12. R.P. REED, P.T. PURTSCHER, N.J. SIMON, J.D. MCCOLSKEY, R.P. WALSH, J.R. BERGER, E.S. DREXLER, AND R.L. SANTOYO, "ALUMINUM ALLOYS FOR ALS CRYOGENIC TANKS: COMPARATIVE MEASUREMENTS OF CRYOGENIC MECHANICAL PROPERTIES OF AL-LI ALLOYS AND ALLOY 2219," REPORT PL-TR--91-3073, PHILLIPS LABORATORY, PROPULSION DIRECTORATE, AIR FORCE SYSTEMS COMMAND, EDWARDS AIR FORCE BASE, 1991 13. C.N. MCCOWAN, T.A. SIEWERT, AND R.L. TOBLER, TENSILE AND FRACTURE PROPERTIES OF AN FE-18CR-20NI-5MN-0.16N FULLY AUSTENITIC WELD METAL AT 4 K, J. ENG. MATER. TECHNOL., VOL 108, 1986 14. T.A. SIEWERT AND C.N. MCCOWAN, THE FRACTURE TOUGHNESS OF 25CR-22NI-4MN-2MO STAINLESS STEEL WELDS AT 4 K, "MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES--XII," NISTIR 3931, NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY, 1990 15. A.O. KLUKEN, C.N. MCCOWAN, AND T.A. SIEWERT, EFFECT OF INCLUSION VOLUME FRACTION AND SIZE DISTRIBUTION ON THE CRYOGENIC TOUGHNESS OF AUSTENITIC STAINLESS STEEL WELD METALS, SUBMITTED FOR PUBLICATION IN MICROSTRUCTURAL SCIENCE, VOL 19, ASM INTERNATIONAL/INTERNATIONAL METALLOGRAPHIC SOCIETY, 1993 16. Y.W. CHENG, H.I. MCHENRY, P.N. LI, T. INOUE, AND T. OGAWA, FRACTURE TOUGHNESS OF 25MN AUSTENITIC STEEL WELDMENTS AT 4 K, ADV. CRYOGENIC ENG. (MATER.), VOL 30, 1983 17. S. TONE, M. HIROMATSU, J. NUMATA, T. HORIUCHI, H. NAKAJIMA, AND S. SHIMAMOTO, CRYOGENIC PROPERTIES OF ELECTRON-BEAM WELDED JOINTS IN A 22MN-13CR-5NI-0.22N AUSTENITIC STAINLESS STEEL, ADV. CRYOGENIC ENG. (MATER.), VOL 32, 1985 18. R.H. ESPY, WELDABILITY OF 21-6-9 STAINLESS STEEL (ARMCO NITRONIC STAINLESS STEEL), "MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES--I, " NBSIR 78-884, NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY, 1978 19. D.J. ALEXANDER AND G.M. GOODWIN, THICK-SECTION WELDMENTS IN 21-6-9 AND 316LN STAINLESS STEEL FOR FUSION ENERGY APPLICATIONS, SUBMITTED TO J. NUCL. MATER., 1992 20. T.A. SIEWERT, PREDICTING THE TOUGHNESS OF SMA AUSTENITIC STAINLESS STEEL WELDS AT 77 K, WELD. J., VOL 65, 1986 21. C.N. MCCOWAN AND T.A. SIEWERT, FRACTURE TOUGHNESS OF 316L STAINLESS STEEL WELDS WITH VARYING INCLUSION CONTENTS AT 4 K, ADV. CRYOGENIC ENG. (MATER.), VOL 34, 1989 22. J.H. KIM, B.W. OH, J.G. YOUN, G.-W. BAHNG, AND H.-M. LEE, EFFECT OF OXYGEN CONTENT ON CRYOGENIC TOUGHNESS OF AUSTENITIC STAINLESS STEEL WELD METAL, ADV. CRYOGENIC ENG. (MATER.), VOL 34, 1989 23. N. YAMAGAMI, Y. KOHSAKA, AND C. OUCHI, MECHANICAL PROPERTIES OF WELDED JOINTS IN MN-CR AND NI-CR STAINLESS STEELS AT 4 K, ADV. CRYOGENIC ENG. (MATER.), VOL 34, 1989 24. C.N. MCCOWAN AND T.A. SIEWERT, INCLUSIONS AND FRACTURE TOUGHNESS IN STAINLESS STEEL WELDS AT 4 K, "MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES--XI," NBSIR 88-3082, NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY, 1988 25. T.A. SIEWERT, C.N. MCCOWAN, AND D.P. VIGLIOTTI, CRYOGENIC MATERIAL PROPERTIES OF STAINLESS STEEL TUBE-TO-FLANGE WELDS, CRYOGENICS, VOL 30, 1990

Welding for Cryogenic Service

T.A. Siewert and C.N. McCowan, National Institute of Standards and Technology

References

1. MECHANICAL, THERMAL, ELECTRICAL, AND MAGNETIC PROPERTIES OF STRUCTURAL MATERIALS, HANDBOOK ON MATERIALS FOR SUPERCONDUCTING MACHINERY, 1ST ED., MCIC-HB-04, METALS AND CERAMICS INFORMATION CENTER, BATTELLE COLUMBUS LABORATORIES, JAN 1977 2. "MATERIALS FOR SUPERCONDUCTING MAGNET SYSTEMS: MECHANICAL AND PHYSICAL PROPERTIES DATA AT LOW TEMPERATURES," BELFOUR STULEN DIVISION, TRAVERSE CITY, MI, 1979 3. T. HORIUCHI, R. OGAWA, AND M. SHIMADA, CRYOGENIC FE-MN AUSTENITIC STEELS, ADV. CRYOGENIC ENG. (MATER.), VOL 32, 1985 4. R.L. TOBLER, T.A. SIEWERT, AND H.I. MCHENRY, STRENGTH-TOUGHNESS RELATIONSHIP OF AUSTENITIC STAINLESS STEEL WELDS AT 4 K, CRYOGENICS, VOL 26, 1986 5. C.N. MCCOWAN AND T.A. SIEWERT, INFLUENCE OF MOLYBDENUM ON THE STRENGTH AND TOUGHNESS OF STAINLESS STEEL WELDS FOR CRYOGENIC SERVICE, "MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES--X," NBSIR 87-3067, NATIONAL BUREAU OF STANDARDS, 1987 6. N.J. SIMON AND R.P. REED, STRENGTH AND TOUGHNESS OF AISI 304 AND 316 AT 4 K, "MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES--X," NBSIR 86-3050, NATIONAL BUREAU OF STANDARDS, 1986 7. R.L. TOBLER, LOW TEMPERATURE FRACTURE BEHAVIOR OF A TI-6AL-4V ALLOY AND ITS ELECTRON BEAM WELDS, STP 651, ASTM, 1978 8. R.L. TOBLER, R.P. REED, AND R.E. SCHRAMM, CRYOGENIC TENSILE, FATIGUE, AND FRACTURE PARAMETERS FOR A SOLUTION-ANNEALED 18 PERCENT NICKEL MARAGING STEEL, J. ENG. MATER. TECHNOL., VOL 100, APRIL 1978, P 189-194 9. J.A. WAGNER, CORRELATION OF MECHANICAL PROPERTIES WITH METALLURGICAL STRUCTURE FOR 18NI 200 GRADE MARAGING STEEL AT ROOM AND CRYOGENIC TEMPERATURES, CRYOGENICS, VOL 31, SEPT 1991 10. R.D. STOUT AND S.J. WIERSMA, FRACTURE TOUGHNESS OF MODERN 9% NICKEL CRYOGENIC STEELS, ADV. CRYOGENIC ENG. (MATER.), VOL 32, 1985 11. J. GLAZER, S.L. VERASCONI, E.N.C. DALDER, W. YU, R.A. EMIGH, R.O. RICHIE, AND J.W. MORRIS, JR., CRYOGENIC MECHANICAL PROPERTIES OF AL-CU-LI-ZR ALLOY 2090, ADV. CRYOGENIC ENG. (MATER.), VOL 32, 1985 12. R.P. REED, P.T. PURTSCHER, N.J. SIMON, J.D. MCCOLSKEY, R.P. WALSH, J.R. BERGER, E.S. DREXLER, AND R.L. SANTOYO, "ALUMINUM ALLOYS FOR ALS CRYOGENIC TANKS: COMPARATIVE MEASUREMENTS OF CRYOGENIC MECHANICAL PROPERTIES OF AL-LI ALLOYS AND ALLOY 2219," REPORT PL-TR--91-3073, PHILLIPS LABORATORY, PROPULSION DIRECTORATE, AIR FORCE SYSTEMS COMMAND, EDWARDS AIR FORCE BASE, 1991 13. C.N. MCCOWAN, T.A. SIEWERT, AND R.L. TOBLER, TENSILE AND FRACTURE PROPERTIES OF AN FE-18CR-20NI-5MN-0.16N FULLY AUSTENITIC WELD METAL AT 4 K, J. ENG. MATER. TECHNOL., VOL 108, 1986 14. T.A. SIEWERT AND C.N. MCCOWAN, THE FRACTURE TOUGHNESS OF 25CR-22NI-4MN-2MO STAINLESS STEEL WELDS AT 4 K, "MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES--XII," NISTIR 3931, NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY, 1990 15. A.O. KLUKEN, C.N. MCCOWAN, AND T.A. SIEWERT, EFFECT OF INCLUSION VOLUME FRACTION AND SIZE DISTRIBUTION ON THE CRYOGENIC TOUGHNESS OF AUSTENITIC

STAINLESS STEEL WELD METALS, SUBMITTED FOR PUBLICATION IN MICROSTRUCTURAL SCIENCE, VOL 19, ASM INTERNATIONAL/INTERNATIONAL METALLOGRAPHIC SOCIETY, 1993 16. Y.W. CHENG, H.I. MCHENRY, P.N. LI, T. INOUE, AND T. OGAWA, FRACTURE TOUGHNESS OF 25MN AUSTENITIC STEEL WELDMENTS AT 4 K, ADV. CRYOGENIC ENG. (MATER.), VOL 30, 1983 17. S. TONE, M. HIROMATSU, J. NUMATA, T. HORIUCHI, H. NAKAJIMA, AND S. SHIMAMOTO, CRYOGENIC PROPERTIES OF ELECTRON-BEAM WELDED JOINTS IN A 22MN-13CR-5NI-0.22N AUSTENITIC STAINLESS STEEL, ADV. CRYOGENIC ENG. (MATER.), VOL 32, 1985 18. R.H. ESPY, WELDABILITY OF 21-6-9 STAINLESS STEEL (ARMCO NITRONIC STAINLESS STEEL), "MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES--I, " NBSIR 78-884, NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY, 1978 19. D.J. ALEXANDER AND G.M. GOODWIN, THICK-SECTION WELDMENTS IN 21-6-9 AND 316LN STAINLESS STEEL FOR FUSION ENERGY APPLICATIONS, SUBMITTED TO J. NUCL. MATER., 1992 20. T.A. SIEWERT, PREDICTING THE TOUGHNESS OF SMA AUSTENITIC STAINLESS STEEL WELDS AT 77 K, WELD. J., VOL 65, 1986 21. C.N. MCCOWAN AND T.A. SIEWERT, FRACTURE TOUGHNESS OF 316L STAINLESS STEEL WELDS WITH VARYING INCLUSION CONTENTS AT 4 K, ADV. CRYOGENIC ENG. (MATER.), VOL 34, 1989 22. J.H. KIM, B.W. OH, J.G. YOUN, G.-W. BAHNG, AND H.-M. LEE, EFFECT OF OXYGEN CONTENT ON CRYOGENIC TOUGHNESS OF AUSTENITIC STAINLESS STEEL WELD METAL, ADV. CRYOGENIC ENG. (MATER.), VOL 34, 1989 23. N. YAMAGAMI, Y. KOHSAKA, AND C. OUCHI, MECHANICAL PROPERTIES OF WELDED JOINTS IN MN-CR AND NI-CR STAINLESS STEELS AT 4 K, ADV. CRYOGENIC ENG. (MATER.), VOL 34, 1989 24. C.N. MCCOWAN AND T.A. SIEWERT, INCLUSIONS AND FRACTURE TOUGHNESS IN STAINLESS STEEL WELDS AT 4 K, "MATERIALS STUDIES FOR MAGNETIC FUSION ENERGY APPLICATIONS AT LOW TEMPERATURES--XI," NBSIR 88-3082, NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY, 1988 25. T.A. SIEWERT, C.N. MCCOWAN, AND D.P. VIGLIOTTI, CRYOGENIC MATERIAL PROPERTIES OF STAINLESS STEEL TUBE-TO-FLANGE WELDS, CRYOGENICS, VOL 30, 1990 Welding for Cryogenic Service T.A. Siewert and C.N. McCowan, National Institute of Standards and Technology

Selected References

• T. OGAWA AND T. KOSEKI, WELDABILITY OF NEWLY DEVELOPED AUSTENITIC ALLOYS FOR CRYOGENIC SERVICE: PART 2--HIGH-NITROGEN STAINLESS WELD METAL, WELD. J., NOV 1987 • R.P. REED, AUSTENITIC STAINLESS STEELS WITH EMPHASIS ON STRENGTHENING AT LOW TEMPERATURES, ALLOYING, J.L. WALTER, M.R. JACKSON, AND C.T. SIMS, ED., ASM INTERNATIONAL, 1988 • R.L. TOBLER, R.P. MIKESELL, AND R.P. REED, CRYOGENIC EFFECTS ON THE FRACTURE MECHANICS PARAMETERS OF FERRITIC NICKEL ALLOY STEELS, STP 677, ASTM, 1979 • K.A. YUSHCHENKO, PROGRESS IN CRYOSTRUCTURAL MATERIALS AND THEIR WELDING IN

THE USSR, ADV. CRYOGENIC ENG. (MATER.), VOL 30, 1983 Welding in Space and Low-Gravity Environments Milo Nance, Martin Marietta Corporation; Jerald E. Jones, Colorado School of Mines

Introduction WELDING AS AN ASSEMBLY PROCESS has become increasingly more attractive to designers of space structures for the same reasons welding is popular on earth. Welding applications in space are similar to those on earth in that they both involve the need to assemble structural components that will exhibit sufficient strength, endurance, and reliability during their service lives. In addition, ease of repair of welded structures is another consideration that makes various welding methods viable for space applications. Typical space welding applications generally fall into three categories: • • •

ON-ORBIT CONSTRUCTION OF LARGE STRUCTURES REPAIR WELDING ON-ORBIT AND FOR EXPLORATION FLIGHTS BEYOND THE VICINITY OF THE EARTH WELDING AND REPAIR OF STRUCTURES ON A LUNAR BASE

This article will review a variety of applications for welding in space and low-gravity environments, describe the unique aspects of the space environment, compare applicable welding processes, and examine the metallurgy of low-gravity welds. In addition, steps that must be taken to ensure the continued development of welding technology in space are discussed. Welding in Space and Low-Gravity Environments Milo Nance, Martin Marietta Corporation; Jerald E. Jones, Colorado School of Mines

Examples of Application Welded joints have the potential to reduce the cost of fabrication and repair on-orbit. For example, large mechanical joints are planned for use on the U.S. Space Station Freedom. Each of these joints weighs approximately 1.36 kg (3 lb). When one considers that the cost for low earth orbit missions is several thousand dollars per pound of payload, it is easy to see how the replacement of these joints with a welded design would save significant amounts of weight and offer the potential to save large amounts of capital. There are other applications for welding technology in space. These involve the modification and repair of structures and components in orbit. It has been suggested that structural material could be scavenged for reuse in orbit. An example of this is the Space Shuttle external tank (Fig. 1). On each shuttle mission the external tank is brought nearly into orbit. With minimal loss of payload capability, these tanks could be brought into orbit allowing the use of the tank structure for experimental platforms or for raw stock material. Cutting and joining technologies could enable the reuse of this valuable resource.

FIG. 1 SPACE SHUTTLE EXTERNAL FUEL TANK

Another application is the repair of major structures. Damage to space structures can result from many causes, including accidental misuse. However, one common cause of damage is "space debris," or miscellaneous small and large pieces of disintegrated satellites, material jettisoned from spacecraft, and other miscellaneous material. This debris is travelling at very high speeds, and may not be in the same orbital path and direction as a spacecraft. A resulting high-speed collision, even with a small piece of material, can cause significant damage. The amount of space debris in low earth orbit is increasing. Figure 2 shows the increase in concentration of debris in orbit from data collected by the Long Duration Exposure Facility (LDEF). The Space Shuttle has already been hit by at least one particle of debris, resulting in a damaged windshield, and there is speculation of impact damage to satellites.

FIG. 2 DIAGRAM SHOWING THE INCREASE IN THE CONCENTRATION OF "SPACE DEBRIS" ENCOUNTERED IN ORBIT. COURTESY OF D.J. KESSLER, NASA JOHNSON SPACE CENTER

On-orbit repair of high-value, long-duration space-based assets has proven to be more effective than to return them to earth. On-orbit welding is a logical technology to assist in the repair of systems such as a space station. In other cases, such as a manned mission to Mars or the vicinity of the moon, it is impractical or impossible to return damaged structures to earth. The Soviet space program already has used welding to repair Soyuz space station tubing. If a lunar base is established, structures will need to be economically constructed. Welding offers the potential for fast and cost-effective fabrication. In addition, repair of structures and equipment damaged by use and/or meteor collision will be necessary. Lunar dust contains many abrasive components, and wear of equipment will be inevitable. The cost of equipment replacement will probably be prohibitive, while repair by weld hardfacing and surfacing is a very costeffective technology that is employed in earth-based applications subjected to wear. See the article "Hardfacing, Weld Cladding, and Dissimilar Metal Joining" in this Volume. Both the United States and the former Soviet Union have carried out welding experiments in space. The United States used Skylab to weld 2219-T87 aluminum alloy, type 304 stainless steel, and pure tantalum samples. The Soviet program used the Soyuz to perform its experiments and still has its space welding technology available. Both efforts involved electron-beam welding. The findings from these experiments are discussed later. Welding in Space and Low-Gravity Environments Milo Nance, Martin Marietta Corporation; Jerald E. Jones, Colorado School of Mines

Space Welding Environment

Space presents a unique welding environment and many technological challenges. The most obvious and dramatic factors are microgravity and high vacuum. Microgravity in earth orbit occurs due to the centrifugal force of the orbiting craft that counters the effects of gravity, making the spacecraft and all orbiting objects, including those within the spacecraft, in continual free fall until a force acts upon them. Typical gravitation accelerations are on the order of 10-6 G (acceleration due to gravity). High vacuum is present around the spacecraft. Typical levels in low earth orbit are 1 × 10-4 to 1 × 10-16 Pa. This vacuum aids some welding processes, such as electron beam. Since virtually no atmosphere is present, there is no attenuation caused by scatter of the electron beam. In addition, for all welding processes there is no need to protect the metal from oxidation or other atmospheric contaminants with a vacuum chamber, flux, or an inert gas atmosphere. The space environment also offers several technological challenges. One of the problems is the lack of a high-capacity electrical power supply. Electrical power generation is currently produced by one of three methods: batteries, fuel cells (power generation from the combination of oxygen and hydrogen to form water), and direct solar conversion. All three of these methods produce a relative small amount of power compared to the requirements of the typical earth-based welding system. Typically, power is limited to 4.5 kW for a period of a few minutes on board the Space Shuttle. This power constraint will be present for the foreseeable future. Larger-capacity generation systems based on thermal-nuclear generators are being developed, but have not yet been used in service with manned spacecraft. A second problem is the lack of nondestructive testing methods qualified for space use. Many of the traditional methods of inspection will not operate well in a space environment. For example, there is no method of dye penetrant inspection developed to date that can be applied in a vacuum. Methods such as radiography and eddy current inspection may be feasible, but no space-qualified equipment has yet been developed. Even if such equipment was available, the competition for weight and volume onboard a spacecraft might prevent its usage. The application of intelligent automation techniques that allow the production of high-quality welds with a high degree of reliability appears to be the most attractive alternative with regard to the available space and weight limitations (see the discussion of the future of space welding later in this article). In order to produce such welds, a clear and precise knowledge of the mechanisms involved in space welding must be developed as well as the equipment necessary to carry out such welds. There have been several advances made in the necessary sensor and control technologies for earth-based welding that can be used in space welding. Welding in Space and Low-Gravity Environments Milo Nance, Martin Marietta Corporation; Jerald E. Jones, Colorado School of Mines

Welding Processes The space environment has constraints that limit the application of certain conventional welding processes. For example, submerged arc welding (SAW) generally depends on gravity to feed the granulated flux and hold it in place on the weld; this is not possible in low-gravity environments. The space environment also affects some of the basic physical processes involved in welding, making welding processes operate differently in space than in earth-based environments. In addition, the weight limitations of orbital payloads on spacecraft will impose constraints on the processes that can be feasibly used in space. The three processes that appear to have good potential for space application are electron-beam welding (EBW), laserbeam welding (LBW), and gas-tungsten arc welding (GTAW). All three processes are often used for autogenous welding, thereby eliminating the need for transport of high-weight consumables. These processes are also amenable to welding in vacuum (with certain modifications to the GTAW process as discussed below). Also, all of these processes have relatively low power requirements when used for thin-section materials typical of spacecraft and space structures, making these processes most useful in space, where all utilities, including electricity, are limited. Electron-Beam Welding (EBW) The EBW process consists of an electron-beam gun that generates a high-energy stream of electrons. When directed through a series of focusing electromagnets, this high-energy stream has sufficient power density to melt metals. The

EBW process is described in detail in the articles "Electron-Beam Welding" and "Procedure Development and Practice Considerations for Electron-Beam Welding" in this Volume. In earth-based environments, EBW generally requires a vacuum since the electrons scatter in a gas atmosphere. In space, however, the natural vacuum environment is ideally suited for the EBW process. Both the United States and the former Soviet Union have experimented with EBW and melting in space. The M512 metals melting experiments using an electron-beam gun, which were conducted by the United States in the Skylab in 1972, are described later in this article. In this case, the operations were performed in a vacuum chamber that was inside the spacecraft, but vented to the outside. At the time of the experiments, there was some concern as to the quality of the vacuum achievable. Often, certain gases and liquids are jettisoned from spacecraft that remain in the environment in the local vicinity of the vehicle and can produce contamination of the vacuum. Advantages. The electron-beam gun is, perhaps, the most efficient of the processes proposed for on-orbit welding. Due

to the power limitations of spacecraft, this process is the most capable for welding of thicker-section materials. For autogenous welding, no consumables are required, thereby eliminating the cost of including them in the spacecraft payload. Limitations. The impingement of electrons onto a metal surface causes the emission of x-rays. The former Soviet Union experimented with a manual electron-beam welding gun for space application. However, x-ray shielding capacity is proportional to the mass of the shield; therefore, shielding of workers may be difficult due to the high cost of payload weight on spacecraft. Automated or robotic welding may be recommended due to the radiation exposure to workers, particularly for extended periods of welding activity. Safety Considerations. Grounding of the part to be welded is very important. Without proper grounding, the part being welded will acquire a net negative charge from the absorbed electrons. This will cause the beam to be deflected and the weld will not form properly. In addition, if the part becomes charged, then it will discharge to anything that it comes in contact with. Such a discharge could be very dangerous for a worker, or could damage electronic equipment such as computers or communications equipment onboard the spacecraft. Although the inherent repulsion of the like-charged electrons will cause beam spread, the beam will remain a fairly concentrated heat source for some distance from the electron-beam gun. Care will be required, both for manual as well as automatic or robotic welds, to ensure that the gun is not activated except when directed at properly prepared components to be welded. Significant damage could occur to workers, to component parts, or to the spacecraft itself if the gun were activated when directed improperly.

Laser-Beam Welding The LBW process consists of a laser and a series of lenses or mirrors that direct and focus the beam at the component to be welded. The process is described in detail in the articles "Laser-Beam Welding" and "Procedure Development and Practice Considerations for Laser-Beam Welding" in this Volume. While the LBW process, unlike the EBW process, does not require a vacuum, the inherent vacuum environment in space will help to prevent oxidation and other contamination of the weld. There are two types of lasers used for welding, gas (CO2) and solid state (Nd:YAG). While the solid-state lasers are relatively low power, this power level is adequate for thin-section components that are likely to be welded in space. Advantages. As with EBW, the LBW process provides a concentrated heat source that can quickly melt through thin

materials without heating large sections of the workpiece. Thus, the necessary power is less than that required for many other processes. However, the laser system is inherently inefficient in the conversion of energy so the input power requirements will be much greater than for an equivalent output heating capacity when compared to EBW. Also, for autogenous welding, no consumables are necessary for LBW, so the payload cost will be low. In addition, the LBW process, unlike EBW, can be used to weld electrically nonconducting materials, including ceramics, glasses, and polymers. Limitations. The inefficiency of the laser system is due to a significant amount of heat generated in the lasing process

itself. This heat will need to be dissipated through the spacecraft radiators or using separate radiators for the laser system. Because heat radiation from low-temperature systems is fairly slow, this may require large radiators to be installed for the laser system; thus the cost and payload weight of the system is increased. Secondly, due to the power input and heat dissipation requirements, it is difficult to produce a laser system for portable application. Recent work with Nd:YAG lasers and fiber optics has resulted in optical fiber delivery of the laser beam at a location remote to the laser system.

However, the laser system generally should be approximately 10 to 30 m (30 to 100 ft) from the LBW torch. Also, great care must be taken to avoid any damage to the cable containing the optical fiber. Safety Considerations. A laser beam will not spread and dissipate with distance as will an electron beam. Consequently, a misdirected laser beam has the potential to cause damage, even over large distances to workers, components, and to the spacecraft. The laser light itself is dangerous to workers if reflected into the eyes. Proper shielding for persons working around laser welding equipment is necessary.

Gas-Tungsten Arc Welding The GTAW process consists of a tungsten electrode connected to one lead of a power source, the other lead of which is connected to the workpiece. When the power supply is activated, current flows to the tungsten electrode through an arc plasma to the workpiece, and returns to the power supply to complete the circuit. The arc plasma is the primary source of heat as electrons move from the metal surface, are accelerated through the plasma column, and impinge on the metal surface at the opposite side of the arc. This welding process is described in the article "Gas-Tungsten Arc Welding" in this Volume. The arc plasma consists of a superheated gas of ions and electrons. To shield the weld from atmospheric contaminants and to provide a stable plasma, inert gas is used. Typically argon, helium, or a mixture of these is chosen for GTAW. Because of the vacuum in the space welding environment, it is not necessary to shield the weld from contaminants; however, a small amount of gaseous material is necessary to sustain a stable arc plasma. The GTAW torch to be used for space welding will be modified from the torch used for earth-based welding. Through the use of a hollow tungsten electrode, the small amount of gas needed to sustain the arc plasma can be provided. A non-reactive gas or gas mixture will be necessary to avoid contamination of the weld. This modified GTAW process is similar to the earth-based PAW process. Advantages. The equipment used for GTAW is simpler and much less expensive to produce and maintain in comparison

to EBW and LBW. Secondly, the arc plasma only exists as a part of the overall welding power circuit, therefore it is nearly impossible to direct the GTAW arc toward a component, worker, or spacecraft that is not connected to this power circuit. Thirdly, as the length of the plasma increases, the voltage required to initiate and maintain the arc increases. By limiting the available voltage in the welding power circuit, the distance that the arc can strike is limited, and the possible damage resulting from a misdirected arc is limited to near-vicinity (25 to 50 mm, or 1 to 2 in.) objects. Limitations. The GTAW arc plasma requires a constant flow of gas. This gas supply, in the form of fairly heavy

compressed gas cylinders, adds to the cost and weight of the spacecraft payload. In addition, this supply will need to be constantly replenished as welding progresses. If resupply is not available, the process will not be available for use. With both LBW and EBW, the welding process would always be available regardless of a resupplied source of welding gas. The arc plasma is a much less concentrated heat source in comparison to EBW and LBW. This results in the need for a higher power feed from the spacecraft to produce a weld in the same thickness material. Safety Considerations. The primary concern with arc welding is grounding and insulation for workers. Should the

worker become part of the welding circuit, the power supply will generally be capable of delivering several kilowatts through the worker at a current level that can cause permanent disability and/or death. Secondly, in order to ionize the inert gas plasma and initiate the arc, a high voltage is usually applied. This is accomplished through the use of a high voltage pulse, or a superimposed high-frequency high-peak voltage signal. Either of these generates a significant amount of electromagnetic interference (EMI) radiation, which can result in damage or shutdown of such equipment as computers and other advanced electronics onboard the spacecraft. Other Joining Processes With a few exceptions (for example, SAW), virtually all welding processes could be used in space. The EBW, LBW, and GTAW generally have been based on weight limitations or on the inherent aspects of the space environment (for example, vacuum for EBW). The materials likely to be found in space that would be welded are typically welded with EBW, LBW, or GTAW in the earth-based environment. These include aluminum and titanium and their alloys. In addition, these processes are well understood, commonly used, and very controllable for thin-section aerospace applications.

Another process that takes advantage of the space environment is solar heat welding. Solar radiation in space is unshielded by the atmosphere. A spacecraft in low earth orbit typically receives 1.4 kW/m2 (0.13 kW/ft2) in daylight conditions. This energy can be taken advantage of by focusing the solar radiation on the workpiece through the use of a large collecting mirror and focusing lenses. Because the optics of the mirror need not be perfect, the cost would be reasonable and the system would not be a drain on the limited power capacity of the spacecraft. Other joining processes that could be utilized include mechanical fastening, solid-state bonding, brazing, and adhesive bonding. However, each of these has inherent limitations that tend to make the three welding processes discussed above more attractive. For example, it may be difficult to use adhesives because they require curing by the application of heat and/or the use of hardening (curing) agents. As a result, adhesives may outgas and contaminate the local environment with volatile components. Mechanical fastening has been used extensively for some joining applications in space. However, mechanical fastening is generally not capable of producing pressurized seals, and mechanically fastened joints are more susceptible to external mechanical damage than are welds. Welding in Space and Low-Gravity Environments Milo Nance, Martin Marietta Corporation; Jerald E. Jones, Colorado School of Mines

Metallurgy of Low-Gravity Welds Influences of Microgravity. The microgravity environment that exists in space applications can have significant effect on the metallurgy of welds. Two important effects on the microstructure of a weld are cooling and fluid flow within the weld pool. In addition, certain other effects can be evident, including the deformation of the weld pool that occurs due to absence of gravity in the space environment. In the latter case, large weld pools in through-penetration or out-of-position welds will not droop or run, but instead their shape will be determined by surface energy and other forces.

Both the cooling rate and the liquid pool stirring behavior are influenced by gravity in earth-based welds. In space, the cooling rate will be different because convective flow is completely suppressed, and this affects the stirring forces in the pool. The principal heat loss from a weld is conduction into the surrounding solid metal. Because there is a significant temperature difference between the weld pool and its surroundings during the time the weld pool is at an elevated temperature, radiative cooling is also important. Generally, convective cooling from the weld surface is only a secondary cooling mechanism until the weld has dropped to a temperature at which little other cooling is occurring and further metallurgical transformation is unlikely to occur. Both the conduction of heat into the surrounding metal and the radiation of heat from the weld pool are directly affected by the local surface temperatures of the weld pool. This includes the exposed surfaces as well as the interface between the liquid pool and the solid base metal. The conductive heat flow is governed by:

Q = K∇ T where ∇ T is the temperature difference vector between the liquid metal and the solid metal at the interface and q is the heat flux vector. This same equation governs the loss of heat by convection from the exposed surface of the weld. However, the value of k is significantly less for heat flow across an interface into a gaseous medium than into a solid metal medium, and the k value for the convective heat transfer in a gas is much less still. The equation for radiative heat flow is:

Q = Eσ( ∇ T)4 where e is the emissivity of the emitting surface and σ is the Steffan-Boltzmann constant. Because the ∇ T term is raised to the fourth power when the temperature difference vector is large, radiation heat transfer is a significant heat loss mechanism for the weld. It should be noted that the radiation of heat from the weld as well as the weld heat source may cause surrounding surfaces to heat, particularly if the weld is made in a vacuum such as outside the spacecraft

environment. The result might be damage to exposed components that are heat sensitive if appropriate shielding is not provided. If convective pool stirring is suppressed, the temperature difference at the liquid-solid interface decreases. The liquid at that surface cools by conduction and the flow of hot liquid from the center of the pool is reduced. The result is a reduced cooling rate for the weld. This reduced cooling rate should tend to cause the grain size in the weld to increase. The opposite effect has, however, been observed and reported by investigators of space-based weld experiments. The absence of gravitational forces also suppresses any buoyancy effects. Because of buoyancy forces, small particles of contaminants in earth-based welds will tend to either float to the weld surface or sink to the weld bottom depending on their density relative to the weld pool liquid. Surface energy effects may cause these particles to accumulate near surfaces. In low-gravity environments, it would be expected that these particles would remain suspended in the liquid pool as it solidifies. If these particles are good nucleants for metallurgical transformations in the weld metal, those transformations would tend to be enhanced. The result would be a difference in the microstructure between earth-based and space-based welds. Differences between the distribution of inclusion particles in earth-based versus space-based welds has been observed. The effects of buoyancy and of cooling rate on actual welds were examined based on experiments conducted on Skylab in June of 1972. Some of those effects are discussed below. Russian Welding Program. Space welding has long been a priority for the space program of the former Soviet Union. In 1969 Soyuz-6 performed the first welds in space on AMG6 and DM-20 aluminum alloys using the Vulkan apparatus. This work was carried on by the E.O. Paton Welding Institute of the Ukraine SSR Academy of Sciences. This work was conducted to develop methods of space repair and construction. Soon thereafter, the United States performed space welding experiments onboard the first manned Skylab Mission in June of 1973. Russian welding efforts have continued, and a portable EBW system has been developed. This hand-held EBW system is currently available for use on aircraft on the Mir space station. United States M512 Experiments. The M551 EBW experiment was conducted in 1973 by astronaut Pete Conrad

inside the multiple docking adapter of Skylab. Figures 3 and 4 show the Skylab spacecraft and the M512 material melting apparatus, respectively. The objectives of the M551 experiment, as described by McKannan and Poorman (Ref 1), were to demonstrate the feasibility of EBW in space and to investigate the differences caused by microgravity. The three materials tested included 2219T87 aluminum alloy, 304 type stainless steel, and tantalum. The aluminum and stainless steel alloys were chosen due to their commonality of use for aerospace applications; the tantalum was chosen as a nearly pure material with no complex phase transformations that would provide information on grain size and other physical and mechanical influences by microgravity. The density, melting point, thermal conductivity, and composition of these base metals are listed in Table 1. In this experiment, an electron beam was used to cut and also produce a weld bead on a 165 mm (6.5 in.) diam rotating disk that varied in thickness from 0.64 to 6.4 mm (0.025 to 0.250 in.) for the stainless steel and aluminum, and from 0.43 to 1.57 mm (0.017 to 0.062 in.) for the tantalum. Finally, the beam was directed at a stationary point on the disk to make a single dwell weld and tO produce a large molten pool. The electron-beam gun was held a constant distance of 38 mm (1.5 in.) with beam adjusted to 80 mA (70 mA for the tantalum sample) and 20 kV. These samples were then compared to samples welded with a duplicate earthbound apparatus using the same welding parameters. All the samples from this experiment were examined in depth at the Battelle Memorial Institute (Ref 2) and at NASA laboratories. In addition, several other organizations received samples for testing and evaluation.

TABLE 1 PROPERTIES AND COMPOSITIONS OF BASE METALS STUDIED IN THE M512 MELTING EXPERIMENTS ALLOY

2219-T87 ALUMINUM TYPE 304 STAINLESS STEEL TANTALUM

DENSITY, G/CM3

MELTING POINT °C °F

COMPOSITION, WT%

2.83

638

1180

THERMAL CONDUCTIVITY W/M · CAL/CM · S K · °C 251 0.6

8.03

1442

2628

159

0.38

...

0.05

...

19.0

0.5

BAL

...

1.5

0.5

8.5

0.04

0.03

0.75

...

...

...

16.6

2996

5425

54.4

0.13

...

...

0.1

...

...

...

...

...

...

...

...

...

...

BAL

...

...

Al

C

Nb

Cr

Cu

Fe

Mg

Mn

Mo

Ni

P

S

Si

Ta

Ti

Zn

BAL

...

...

...

6.2

0.25

0.02

0.3

...

...

...

...

0.15

...

0.15

0.10

FIG. 3 PHOTOGRAPH OF SKYLAB, WHICH HOUSED THE M512 METALS MELTING EXPERIMENT

FIG. 4 PHOTOGRAPH OF THE M512 MELTING APPARATUS TAKEN AT THE AIR AND SPACE MUSEUM IN WASHINGTON D.C.

The metallurgical study concluded that electron-beam space welding was feasible. However, distinct changes in weld morphology and metallurgy were found. These are reported in Table 2.

TABLE 2 DIFFERENCES IN SPACED-BASED (SKYLAB) AND EARTH-BASED WELD SAMPLES

CHARACTERISTICS OF SKYLAB SAMPLE INCREASED OXIDE ON BACK SIDE OF WELD GRAIN REFINEMENT IN WELD REDUCTION IN SURFACE RIPPLE, BEADING, AND SPATTER INCREASED NUGGET SYMMETRY 304 STAINLESS STEEL GRAIN AND SUBGRAIN REFINEMENT DECREASED BANDING DECREASED WELD POOL SAG REDUCTION IN SURFACE RIPPLE, BEADING, AND SPATTER INCREASED NUGGET SYMMETRY TANTALUM GRAIN REFINEMENT DECREASED WELD POOL SAG INCREASED NUGGET SYMMETRY MATERIAL 2219-T87 ALUMINUM

A subsequent investigation by Nance and Jones at the Colorado School of Mines has shown that both the grain size and the inclusion distribution were affected by the microgravity environment. This study concentrated on the tantalum material. It had been previously hypothesized that the grain size of the weld made in space would be larger due to the reduced cooling rate of the liquid pool. This hypothesis was contradicted by the observation of a 23% reduction in grain size observed in the Skylab tantalum welds in comparison to the earth-based welds. Figures 5(a) and 5(b) compare the grain size of the earth-based and Skylab tantalum welds. A similar grain size reduction was also noted in the aluminum and stainless steel alloys. The mechanism for grain size refinement in these alloys is complex, and it is difficult to draw

direct conclusions. For the pure tantalum samples, there were no solid-state transformations; however, the grain size change may be attributable to an inoculant nucleation mechanism.

FIG. 5 COMPARISON OF (A) GROUND-BASED AND (B) SKYLAB-BASED WELDS IN PURE TANTALUM SAMPLES. NOTE THE LARGER GRAIN SIZE IN THE GROUND-BASED WELD. SOURCE: REF 2

Utilizing scanning electron microscopy (SEM) with energy dispersive x-ray analysis and scanning Auger spectroscopy (SAS), it was determined that oxide inclusions were well distributed throughout the Skylab welds. In the earth-based welds, these oxide particles, possibly due to the influence of gravity-induced buoyant forces, had floated to the top surface of the weld pool. Figure 6(a) is a SEM micrograph of the Skylab tantalum weld. There is a large carbide particle near the center of the SEM photomicrograph, as confirmed by the SAS carbon map shown in Fig. 6(b). Figure 6(c) is the SAS oxygen map, showing only a small accumulation of oxide particles near the surface of the weld, with a well-distributed oxide particle dispersion in the center of the weld. The earth-based weld showed virtually no oxide particles in the middle of the weld and a very heavy accumulation of oxides at the weld surface.

FIG. 6 MATERIALS CHARACTERIZATION TECHNIQUES FOR EXAMINING A TANTALUM WELD CARRIED OUT IN SKYLAB. (A) SEM MICROGRAPH SHOWING A LARGE CARBIDE PARTICLE NEAR THE CENTER OF THE SAMPLE. (B) SCANNING AUGER SPECTROSCOPIC CARBON MAP SHOWING A LARGE CARBIDE PARTICLE NEAR THE CENTER OF THE SAMPLE. (C) SCANNING AUGER SPECTROSCOPIC OXYGEN MAP SHOWING ONLY A SMALL ACCUMULATION OF OXIDE PARTICLES NEAR THE SURFACE OF THE WELD WITH A WELL-DISTRIBUTED OXIDE PARTICLE DISPERSION IN THE CENTER OF THE WELD


1. E.C. MCKANNAN AND R.M. POORMAN, PAPER 35812, GEORGE C. MARSHALL SPACE FLIGHT CENTER, ALABAMA, 1974, P 20 2. R.E. MONROE, "CHARACTERIZATION OF METALS MELTING RISKS: SKYLAB EXPERIMENT M551," NASA CONTRACTOR REPORT, BATTELLE MEMORIAL INSTITUTE, COLUMBUS, OHIO, 4 DEC 1973, P 109

Welding in Space and Low-Gravity Environments Milo Nance, Martin Marietta Corporation; Jerald E. Jones, Colorado School of Mines

Future of Space Welding The most important space-based welding application for long-duration manned spacecraft is repair welding. This technology offers the possibility of decreasing assembly time and decreasing the weight and cost of space structures. However, before welding technology can become commonplace in space, several steps must be taken. These involve the development of processes and controls that allow highly reliable welds to be performed in space. These steps include: •

•

•

DEVELOPMENT OF EQUIPMENT THAT HAS SUFFICIENT SENSOR TECHNOLOGY TO REPORT IN REAL TIME THE QUALITY OF WELDS AND TO FACILITATE REAL-TIME PROCESS CONTROL AND CORRECTION OF WELDS IN PROGRESS CHARACTERIZATION OF THE DIFFERENCES IN SPACE-BASED WELDS AND EQUIVALENT EARTH-BASED WELDS TO PROVIDE A BASIS FOR MODELING AND UNDERSTANDING OF THE DIFFERENCES IN THESE WELDS PROCESS DEVELOPMENT FOR WELDS CONDUCTED IN SPACE

If these three actions are taken, welding can potentially become the preferred method of joining metallic components in space and an important method for the long-term maintenance of structures, including large space-borne structures that are independent of earth. The Need for Automation. Earth-based welding is a difficult task and requires both significant skill as well as time-

consuming and tedious effort. If performed manually, space welding will be, at best, a very difficult task. The problems with space welding are compounded if the welding is to be done outside the spacecraft environment. Extra-vehicular activity (EVA) requires large numbers of ground personnel and considerable time and preparation by the mission specialist. In order for spacecraft crews to have the skill level required to make critical welds, considerable time and additional effort would have to be expended. Such training would be very expensive. Computer intelligent autonomous or semi-autonomous welding systems will probably be the most cost-effective means for space welding, particularly for EVA welding. Intelligent Engineering/Planning Systems. Recently developed systems have demonstrated the possibility of

coupling extensive databases with artificial-intelligence-based software to produce "intelligent" engineering and planning software that can make many routine engineering decisions for welding applications and guide a human user to optimum decisions for welding. These systems utilize expert system and neural network technology to provide accurate weld process models and engineering decision-making capability. This subject is addressed for earth-based welding applications in the article "Intelligent Automation for Joining Technology" in this Volume. Intelligent engineering and planning workstations that are electronically coupled to automatic welding machines will substantially reduce the need for time-consuming work by mission specialists. The planning and development of welding schedules and procedures can be accomplished in a matter of a few minutes on a computer terminal rather than hours of work by hand. By having a complete set of databases and most of the routine engineering completed by the computer, the human user will need only to answer a few questions about the application, and the system will produce a complete welding plan and weld schedule that can be downloaded directly to the automated welding system. Such systems will offer considerable cost savings and manhour reduction while at the same time providing a carefully optimized welding plan and procedure.

Welding in Space and Low-Gravity Environments Milo Nance, Martin Marietta Corporation; Jerald E. Jones, Colorado School of Mines

References

1. E.C. MCKANNAN AND R.M. POORMAN, PAPER 35812, GEORGE C. MARSHALL SPACE FLIGHT CENTER, ALABAMA, 1974, P 20 2. R.E. MONROE, "CHARACTERIZATION OF METALS MELTING RISKS: SKYLAB EXPERIMENT M551," NASA CONTRACTOR REPORT, BATTELLE MEMORIAL INSTITUTE, COLUMBUS, OHIO, 4 DEC 1973, P 109 Joining of Organic-Matrix Composites L.J. Hart-Smith, McDonnell Douglas Aerospace

Introduction SUCCESSFUL ADHESIVE BONDING of organic-matrix composites is dependent on the nature of the adherend surfaces. Factors that influence the integrity and durability of composite-to-composite bonds include wetting of the adherend, cleanliness of the substrate, chemical and physical properties of the adherends, and joint design. In this article, major emphasis is placed on the critical importance of proper surface preparation; both thermoset and thermoplastic composites are covered. The techniques used to analyze and design adhesively bonded composite joints are described in a companion article (Ref 1). Additional information can also be found in Adhesives and Sealants, Volume 3 of the Engineered Materials Handbook published by ASM International.

Reference

1. L.J. HART-SMITH, JOINTS, COMPOSITES, VOL 1, ENGINEERED MATERIALS HANDBOOK, T.J. REINHART, ED., ASM INTERNATIONAL, 1987, P 479-495 Joining of Organic-Matrix Composites L.J. Hart-Smith, McDonnell Douglas Aerospace

Common Problems Encountered The Importance of Proper Processing. It is well known that the adhesive bonding of metal structures requires strict adherence to the process specifications if the glue is to remain stuck for a long service life. It is also apparent that not all approved specifications are adequate. The Primary Adhesively Bonded Structures Technology (PABST) program (Ref 2) was conducted to demonstrate the tremendous benefits to metallic aircraft structures from good adhesive bonding, to associate previous problems in service with processing techniques that should be discontinued, and to validate the use of phosphoric acid anodizing with a corrosion-inhibiting primer as a reliable technique--as durable as the British chromic acid anodizing with Redux adhesive, and easier to use.

Only when the adhesive is stuck securely to the adherends is it possible to develop a bond strength that can be relied upon. Only then is it possible to use rational design procedures in which the joint can be designed for the bond to never fail, or if it does fail, for the failure to be characterized by cohesive failure of the adhesive layer itself. The interfacial failures associated with improper or inadequate processing are both unpredictable and variable--so variable, in fact, that

neither the strength nor fatigue life can be established. This poor processing is the source of the so-called "weak bonds" that are discovered only after the structure has fallen apart. Industrial experience with bonding of both metals and composites has shown that it is very easy to fabricate defective bonds that are strong enough to pass all initial inspections, but that have only a small fraction of the strength to be expected from properly processed bonds. Techniques for creating suitable surface preparations must be followed to achieve the normal situation in which the bonds are so much stronger than the thin members they join that all failures will occur outside the joint. The necessary steps are not difficult, but are often omitted because of a failure to appreciate the need for them. Unfortunately, the importance of surface preparation for the adhesive bonding of fibrous composites is not widely acknowledged. Robert Schliekelmann, the pioneer of Redux bonding at Fokker, was sufficiently concerned about this oversight to make a plea for more attention to this issue (Ref 3). Hart-Smith and co-workers at Douglas Aircraft Company prepared a document on surface preparation of fibrous composites to help the airlines until repair manuals were completed (Ref 4). When this document was reprinted in the Canadair house journal (Ref 5), every experiment had to be repeated to create new photographs and every phenomenon was duplicated in order to verify that the problems discussed in the Douglas Aircraft article really existed. A similar concern was expressed in England, where Parker and Waghorn (Ref 6) reported on a far more comprehensive test program on the effects of surface preparations on adhesive bond strength for carbon-epoxy laminates. They also concluded that composite surfaces must be abraded to achieve strong adhesive-bonded joints. Pocius and Wenz (Ref 7) advocated the use of Scotchbrite pads with embedded abrasive particles as an effective and reliable surface preparation technique for achieving good composite bonds (Ref 7). The form of Scotchbrite giving the best results is of the consistency now used widely in floor-scrubbing pads, which is a little finer than that used in rotary discs for stripping rust from automobiles before repainting. Sources of Defective Bonds. A major source of defective bonding of composite adherends is the omission of a thorough mechanical abrasion of the surface, either by a low-powered grit-blast machine or thorough sanding, following the removal of a peel ply from the surface of the laminate. Grit blasting is shown to lead to the best thermoset bonds, which are enhanced by the use of an adhesive film cured on the surfaces of the part as it is made. Although peel plies can protect against gross contamination, they do not create a surface that is easily wetted by adhesives. Despite the cautions about the many wrong ways to make bonded joints, correct processing is neither difficult nor expensive and results in strong durable structures.

Most, if not all, peel-ply surfaces are inadequate for adhesive bonding. This is a contentious issue because many bonded composite structures are so thin that the laminate or honey-comb core fails first during testing, even when the bond is less than 10% effective. However, that is often not the case in service. It has sometimes been found that when a part damaged in service is disassembled to provide access to abrade the surfaces for a bonded or co-cured repair, it is relatively easy to break the bonds apart without delaminating much of the composite. In such a case, all of the original bond area that was neither detected as defective during initial inspection, nor delaminated as the doublers and stiffeners were removed from the skin, should be regarded as weak (defective) bonds that could not be detected by ultrasonic inspection. This fraction of the bond area may be as high as 90% of the total if the original bonding had not been processed correctly. This condition is usually associated with virtually all of the adhesive on one surface or the other, when assessed locally, although the surface to which the adhesive adhered may vary throughout the total bond area. The local areas in which there was no adhesion are typically quite large and cannot be mistaken for porosity. It is extremely rare to see half the bond completely covering each surface after disassembly; the cohesive strength of the adhesive layer, if it has been processed correctly, is usually so much greater than that of the adherends that trying to pry good bonds apart will result in delaminating the composite, not the adhesive layer. As a point of reference for the tremendous peel strength of most modern adhesives, attempts to salvage machined metal fittings from completed test panels for reuse on later panels in the PABST bonded metal fuselage program were largely fruitless, even when the structure was struck with sledge hammers and cold chisels after chilling the panels with liquid nitrogen. (Some older adhesives have so little peel strength at low temperature that it used to be possible to disassemble bonded metal structures by giving them a sharp rap after chilling them with dry ice.) It would seem to make more sense to heat bonded composite structures above the glass transition temperature for the adhesive, but below that of the parent laminate to facilitate safe disassembly. The disadvantage here would be that the working environment of some 120 °C (250 °F) would be rather unpleasant, unless it could be confined to the immediate vicinity of the area of prying by using a jet of hot air.

There is no history of bonds that were truly good at the time of manufacture deteriorating in service. There have, however, been numerous instances in which the interface of a metal bonded structure has corroded in service because it had not been properly prepared in the first place and was just sound enough to pass the initial inspection. Although poor composite bonds deteriorate by mechanisms other than corrosion, they share with poor metal bonding the characteristic that their short-term static strength is sufficient to conceal their deficiencies. Weak bonds have been notoriously difficult to detect prior to failure when only the parts have been examined; however, they are easy to identify after failure and can be recognized reliably from a straightforward reading of the process monitoring records and an on-site survey of the procedures followed (and ignored) during manufacture. The process monitoring records should always take precedence over successful test coupon results, even when it means rejecting and rebuilding (not reworking) complete assemblies. The reason for the difficulty in bonding of polymer composites is that it is difficult to create strong chemical bonds between the adhesive and a cured (and therefore largely fully crosslinked) polymer matrix. The cured surface needs to be activated and, even then, chemical bonds may be degraded by absorbed water (Ref 8). The most durable bond strengths rely upon mechanical interlocking of the adhesive and the thermoset composite surface, just as with the primer and porous anodized surface on metal adherends. If the surface is not adequately roughened, the glue will not stick, even if the surface is not contaminated. The most common contaminants are industrial-grade cleaning solvents and coupling fluids for ultrasonic inspections. These are particularly difficult to remove from delaminations and areas of porosity. Failure to remove them completely results in a bond with inadequate adhesion to the adherend, just as is the case with a metal-bond surface that had been contaminated before the primer was applied. Effects of Absorbed Moisture. Water is the most widespread and troublesome source of defective bonded joints and

repairs in composite structures. It may be absorbed by the basic composite material before it is ever cured. If it is not removed by drying prior to curing or by venting during cure, it will result in porosity as well as harmful changes in the chemical structure of the polymer. Water may also be absorbed by uncured adhesive films that have been improperly stored or that have not thawed out before being unwrapped after removal from the freezer or that have been subjected to an extended layup time in a humid work environment. Adhesives vary in susceptibility to this problem, but most have a far greater affinity for water than do epoxy prepregs. The consequences are much the same, and just as detrimental. Water may also be absorbed within cured laminates that are to be secondarily bonded either during initial manufacture or as part of an in-service repair. If this water is not removed by slow drying prior to bonding, the water will react adversely with the adhesive and result in defective bonds. Cured epoxies have very low diffusion coefficients, so water is not absorbed very rapidly. However, they cannot be dried rapidly for the same reason. A systematic series of tests to differentiate between the effects of different kinds of prebond moisture has not been carried out to date. However, a weak, porous adhesive that crumbles easily into powder when the fractured surface is handled after the adhesive layer has been exposed, but that has stuck to the adherends, can be associated with too much moisture absorbed by the adhesive before cure. A perfectly sound solid adhesive film that simply did not stick is more likely to be caused by precured laminates to be bonded that were not dried or by condensate on the surface of an otherwise dry adhesive film that was unwrapped too soon after removal from storage in the freezer. However, some film adhesives left out for several days in a humid layup area can absorb so much water that they will not stick even if they are not visibly covered by condensate during layup. Even then, there is a chance that the strength of the fillets at the edges of the bonds will be sufficient to disguise the weakness by preventing the parts from falling apart after completion of the cure. Labor and Myhre (Ref 9) found that if they did not dry the parent laminate first, they could not attach even a single 0.13 mm (0.005 in.) thick carbon-epoxy tape without serrating the ends with pinking shears. With square-cut ends, the single layer would rapidly delaminate under fatigue loads of only about 25% of design limit load. Obviously, however, such weak bonds would have passed all nondestructive inspections at the time the repair was executed. The only way of knowing that the bond would have failed had it been placed in service was recognition that water definitely was present during the cure of the adhesive--water that was added deliberately to make the coupons representative of structures in service that were being quickly repaired without first being dried out slowly and carefully. In extreme cases, with a rapid temperature rise during bonding, the parent laminate may even split apart because of the steam generated by the water that was unable to escape fast enough. Figure 1 shows the dramatic reduction in the hot-wet strength of adhesively bonded joints caused by as little as 0.25% absorbed moisture in the composite laminates before bonding. Less than 20% of the room-temperature strengths were retained. Figure 2 shows equally dramatic losses in the fatigue life under slowcycle testing of bonded aluminum joints as a function of different levels of prebond moisture in the adhesive film prior to bonding.

FIG. 1 EFFECT OF MOISTURE CONTENT ON THE HOT-WET BOND STRENGTH OF CARBON-EPOXY COMPOSITES. SOURCE: REF 10

FIG. 2 EFFECT OF PREBOND MOISTURE ON FATIGUE STRENGTH OF ADHESIVELY BONDED ALUMINUM JOINTS. FILMS WERE PRECONDITIONED FOR 10 DAYS AT ROOM TEMPERATURE. RAAB SPECIFICATIONS, 10.3 MPA (1500 PSI), 4 CYCLES/H, 60 °C (140 °F), 100% RELATIVE HUMIDITY (RH). SOURCE: REF 11

Although the evidence is not yet absolutely conclusive, the epoxy resins may be more tolerant of absorbed moisture than is generally acknowledged if moisture can easily be driven off early in the cure cycle. However, if there are no vents to allow the steam (or solvents) to escape, as is often the case with no-bleed cure cycles, moisture cannot be driven off. There is conclusive evidence that the combination of the presence of moisture and a closed cure environment is

disastrous. However, there have been some instances with metal-bonded structures when water was known to be present in the uncured adhesive, yet the bonds appear to have been good (per service). In each case, the bonding tool permitted easy egress of the steam generated during cure. Companion panels manufactured with a different bonding tool that trapped the moisture in the adhesive from the very same rolls, passed all ultrasonic inspections but were found to be unbonded during assembly. A similar problem arose with the two-stage bonding of a composite spoiler. Although water was present in some form during both bonds, there was adequate venting for the steam to escape while the first skin was bonded to an oversize block of honeycomb core, to stabilize it for subsequent machining of the aerodynamic profile (it is not known if the water was absorbed in the uncured adhesive film, if the Nomex honeycomb core had not been dried properly, or if the precured laminates had not been dried properly). This first bond was so strong as to give no warning of the problems to follow. When the second skin was bonded to the other side of the core, the water was trapped and the second bond had only a small fraction of the strength exhibited by the first bond. When tested, the second bond failed prematurely; it looked porous and weak, with poor filleting at the cell walls, while the first bond showed good filleting, as if the water had never been present. The above examples are included not as an invitation to disregard the concerns about moisture present during the cure, but as encouragement to pursue this subject further to see if there are any consistently reliable alternatives to careful drying prior to bonding, which is obviously time consuming and costly. It is also intended to emphasize the need for venting during cure. There are other known instances in which simple thin laminates have suffered locally from very wavy fibers caused by the lack of an escape path for bubbles of gas created from the solvent used during the process of impregnating the fibers with resin. The bubbles lifted the fibers off the tool surface, pulling the ends in toward the bubbles. Eventually, the bubbles burst, causing the fibers to wrinkle as they were flattened out. There was no mechanism to push the fibers back to where they had been. Difficulties in Detecting Defective Bonds. The most difficult aspect of defective bonds caused by prebond moisture or inadequate surface preparation is that the condition cannot be detected by ultrasonic inspection. It must be avoided by adherence to the correct procedures--which may differ from those specified--during processing because the weak bonds resulting from such errors cannot be repaired, even if they are detected before they fail. The parts must be disassembled, the surfaces cleaned and prepared again and the entire bonding cycle repeated. In general, it is easy to rework improperly manufactured mechanically fastened metallic structures and to restore 100% of the design strength. That is not true for either adhesively bonded metal or composite structures; the rework is difficult and the strength recovered far less than complete. This is why it is so important to do the job properly in the first place.

The cost of accepting defective parts is compounded as the original panels are embedded ever more deeply in the structural assemblies. It is bad enough when the defects are unrecognized because the end-item inspections are ineffective; inspection can exclude a few very bad parts, but cannot ensure the integrity of the great majority of assemblies. It is worse when parts are recognized as defective from a visual inspection, but are accepted anyway on the basis of coupon tests or ultrasonic inspections. Bonding of Thermoplastic Composites. Thermoplastics are far more difficult to bond than are thermoset composites,

but good results have been achieved by chemically etching the surface prior to bonding. The most promising approach for joining thermoplastic composites appears to be fusion bonding, using a compatible surface layer of plastic that melts at a lower temperature than that needed to melt the parent matrix. The ability of typical modified epoxy adhesives to bond chemically to the thermoplastic matrix is even less than for thermoset composites. Also, the thermoplastic surface is so resilient that it is very difficult to abrade sufficiently without breaking too many surface fibers in the process of creating the roughened bond surface. The best bonds made with layers of thermoset adhesives between thermoplastic adherends seem to have a strength of only about 80% of that of the best bonds using the same adhesive between thermoset adherends. One must conclude, therefore, that this loss of strength results from an interfacial bond strength that is weaker than the intrinsic cohesive strength of the adhesive layer. It follows, therefore, that the strength to be attained would be highly variable, as would the service life. Far better success has been achieved by chemically changing the surface to be bonded by oxidation using hot gases, by cold gas plasma treatment, or by corona discharge. The surface energy is increased by each of these techniques, permitting better wetting by the adhesive. However, these processes are rather slow if a large area has to be covered. Success in bonding thermoplastics has also been achieved by fusion, or welding. The best success has been achieved when the original thermoplastic laminates have been fused together with a special surface layer of a different polymer that can be melted at a lower temperature than is needed to melt the base laminates. Direct fusion of thermoplastic laminates often results in considerable distortion, even if a structurally sound bond is achieved. Even making thermoplastic

composites using focused local heat to fuse each ply to the substrate has proved to be a slow process that is not well suited to large panels. Similarly, it is common practice in making thermoset composite aircraft panels to include a surfacing layer of film adhesive co-cured with the parent matrix to both facilitate getting the paint to adhere and to provide a tough gel coat to compensate for the excessively dry surface plies in normal composite laminates. (The weight of this surfacing layer is less than that of the glazing that would otherwise need to be applied, by hand, to avoid pinhole porosity and an unacceptable paint finish.) Curiously, less use is made of this concept for preparing surfaces to be bonded, although an Australian firm has reported great success with it (Ref 12).


2. E.W. THRALL, JR. AND R.W. SHANNON, ED., ADHESIVE BONDING OF ALUMINUM ALLOYS, MARCEL DEKKER, 1985 3. R.J. SCHLIEKELMANN, ADHESIVE BONDING AND COMPOSITES, PROGRESS IN SCIENCE AND ENGINEERING OF COMPOSITES, VOL 1, PROC. 4TH INT. CONF. COMPOSITE MATERIALS (TOKYO), 22-28 OCT 1982, JAPAN SOCIETY FOR COMPOSITE MATERIALS AND THE METALLURGICAL SOCIETY (TMS) OF AIME, T. HAYASHI, K. KAWATA, AND S. UMEKAWA, ED., NORTH-HOLLAND, AMSTERDAM, 1982, P 53-78 4. L.J. HART-SMITH, R.W. OCHSNER, AND R.L. RADECKY, SURFACE PREPARATION OF FIBROUS COMPOSITES FOR ADHESIVE BONDING OR PAINTING, DOUGLAS SERVICE MAGAZINE, FIRST QUARTER 1984, P 12-22 5. L.J. HART-SMITH, R.W. OCHSNER, AND R.L. RADECKY, SURFACE PREPARATION OF FIBROUS COMPOSITES FOR ADHESIVE BONDING OR PAINTING, CANADAIR SERVICE NEWS, VOL 14 (NO. 2), SUMMER 1985, P 2-8 6. B.M. PARKER AND R.M. WAGHORNE, SURFACE PRETREATMENT OF CARBON FIBREREINFORCED COMPOSITES FOR ADHESIVE BONDING, COMPOSITES, VOL 13, JULY 1982, P 280-288 7. A.V. POCIUS AND R.P. WENZ, MECHANICAL SURFACE PREPARATION OF GRAPHITE-EPOXY COMPOSITE FOR ADHESIVE BONDING, SAMPE J., SEPT/OCT 1985, P 50-58 8. A.J. KINLOCH AND G.K.A. KODOKIAN, "THE ADHESIVE BONDING OF THERMOPLASTIC COMPOSITES," EUROPEAN RESEARCH OFFICE OF THE U.S. ARMY, CONTRACT DAJA 45-86-C0037, FOURTH PERIODIC REPORT, IMPERIAL COLLEGE, LONDON, AUG-DEC, 1987 9. J.D. LABOR AND S.H. MYHRE, "LARGE AREA COMPOSITE STRUCTURE REPAIR," NORTHROP AIRCRAFT CORPORATION, USAF TECHNICAL REPORT AFFDL-TR-79-3040, MARCH 1979 10. B.M. PARKER, THE EFFECT OF COMPOSITE PREBOND MOISTURE ON ADHESIVE-BONDED CFRP-CFRP JOINTS, COMPOSITES, VOL 14, JULY 1983, P 226-232 11. E.W. THRALL, JR., PROSPECTS FOR BONDING PRIMARY AIRCRAFT STRUCTURES IN THE 80'S, 25TH NATIONAL SAMPE SYMPOSIUM, VOL 25, SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 6-8 MAY 1980, P 716-727 12. I. MCARTHUR, HAWKER DEHAVILLAND, BANKSTOWN, AUSTRALIA, PRIVATE COMMUNICATION, 1992 Joining of Organic-Matrix Composites L.J. Hart-Smith, McDonnell Douglas Aerospace

Examples of Good and Bad Adhesive Bonds

Properly Processed Bonds. Figure 3 shows the magnified fractured surface of a properly processed adhesive bond,

with continuous coverage of the adherend. With composite adherends, this condition can occur only in test coupons deliberately designed with so little bond area that the bond will fail by shear. It is also necessary to suppress the generation of peel stresses at the ends of the overlap, by keeping the tip thickness of the adherends low. Properly designed structural joints will always fail outside the bond area, and while a coupon test based on such a joint can confirm the adequacy of the design, it cannot confirm that the processing was done properly because there is no fractured bond line to inspect. In properly designed bonded joints, the integrated shear strength of the bond should exceed the load needed to fail the adherends. It is almost impossible not to satisfy this condition for thin adherends, even for improperly processed bonded joints that will fail in service, but are strong enough to pass all the initial inspections. However, for thick adherends, it becomes necessary to use stepped-lap joints to prevent the bond from becoming a weak-link fuse.

FIG. 3 COHESIVE FRACTURE SURFACE IN A PROPERLY PROCESSED COMPOSITE-TO-COMPOSITE ADHESIVE BOND. VISIBLE FIBERS ARE THE RANDOM MAT CARRIER IN THE ADHESIVE FILM. 68×

Adhesion Failures.

*

Figure 4 shows a typical adhesion failure in which the adhesive almost totally failed to adhere to one of the adherends because the surface had not been adequately roughened. Although this photomicrograph is of the surface of the adhesive, not the composite, it clearly shows the imprint of the peel-ply woven surface to which it had not bonded over most of the surface. The comparison between the two modes of failure is dramatic. This failure of the adhesive to bond to a surface arises frequently when an attempt is made to bond directly to the surface that remains after removal of a peel ply, but it also occurs when one tries to bond directly to the other surface of the laminate, which is usually contaminated by release agents used to prevent the part from sticking permanently to the layup tool.

FIG. 4 ADHESION FAILURE SHOWING IMPRINT OF THE PEEL PLY AND FAILURE TO BOND OVER MOST OF THE SURFACE. A SMALL PIECE OF RESIN ADHERED AND IS IDENTIFIED BY THE IMPRINT OF THE FIBERS. 27×

In a different kind of adhesion failure, the adhesive does not remain in one continuous disbonded layer, but remains attached to one of the adherend surfaces in discrete full-thickness fragments. The underlying shiny peel-ply surface is clearly evident in the holes in the adhesive on both surfaces. Roughly half the adhesive is attached to each surface, but

only weakly. This kind of fracture has been observed in uncontaminated bond surfaces that had not been mechanically abraded to make the glue stick. It is the author's belief that this condition, in contrast to that shown in Fig. 4, can occur only in conjunction with small amounts of moisture trapped in the adhesive as it cures. Given that the interfacial bond is uniformly poor, it seems inconceivable that it could have been strong enough to tear a properly cured adhesive film from side to side. However, it is well known that a weak powdery bond results if the adhesive is cured in the presence of moisture. This moisture can either be absorbed by the adhesive film, through improper storage and handling, or by the laminate during service (with even more absorbed moisture, some adhesives will not stick at all). Actually, even the time involved between detail manufacturing and assembly can be sufficient to permit the absorption of sufficient moisture to create a weak powdery adhesive bond, even if the bond surfaces had been properly prepared by thorough abrasion. The laminates need to be stored in sealed containers until bonded or in a dehumidified air-conditioned room. Inspection Methods. None of the structural deficiencies associated with weak or improperly cured adhesive bonds are

detected by conventional ultrasonic nondestructive inspection (NDI), and it is most improbable that they could be detected prior to a gap being opened up by mechanical failure in service. Although the fillets at the ends of the bond overlap will be visibly porous for adhesives cured with trapped prebond moisture, it is usually only the hot-wet lap-shear tests that fail to meet the specified strengths. Consequently, the customary room-temperature tests that follow would not have caught this condition had it been on a real part. Figure 5 shows why visual inspections of adhesive bonds can be more informative and reliable than conventional ultrasonic NDI. More reliable results can be obtained from a visual inspection of the fillet than from either ultrasonic inspections or a defined "adequate" lap-shear test result on a coupon with abnormal fillets. Ironically, the ultrasonic signals are unreliable in the immediate vicinity of the edges of panels-which are the only locations where any significant load is transferred through bonded joints.

FIG. 5 EXAMPLES OF GOOD AND BAD ADHESIVE FILLETS

Note cited in this section

* *IT IS A COMMON MALPRACTICE TO USE THE WORDS "ADHESIVE FAILURES" EXCLUSIVELY TO DESCRIBE ONLY THOSE SITUATIONS IN WHICH THE ADHESIVE ITSELF HAS NOT FAILED, BECAUSE ADHESION IS A NOUN AND NOT AN ADJECTIVE. ACTUAL FRACTURE OF THE ADHESIVE ITSELF IS DESCRIBED AS A COHESIVE FAILURE. THE FORMER FAILURE OCCURS WHEN THE GLUE HAS NOT ADHERED TO ONE OR BOTH ADHERENDS. THE CURRENT MEANING ASCRIBED TO ADHESIVE FAILURES IS THUS VERY MISLEADING, EVEN IF IT IS GRAMMATICALLY CORRECT. IN THIS ARTICLE, THE WORD ADHESION IN PLACE OF ADHESIVE WILL BE USED IN THIS CONTEXT.

Joining of Organic-Matrix Composites L.J. Hart-Smith, McDonnell Douglas Aerospace

Surface Preparation for Thermoset Composite Adherends While not always used in production, light grit blasting, thorough sanding, or some other form of mechanical abrasion are the best surface preparations for both adhesive bonding and painting of thermoset composites. This is confirmed by the test results shown in Fig. 6. As shown in this figure, the best example of sanding was as good as the grit-blasted coupons, but all other treatments were markedly inferior.

FIG. 6 EFFECT OF SURFACE PREPARATION ON THE STRENGTH OF DOUBLE-LAP GRAPHITE/EPOXY-TOALUMINUM ADHESIVE JOINTS (13 MM, OR

1 IN., OVERLAP) 2

Peel-Ply Surfaces. While it is extremely difficult to ensure that a peel-ply surface will be satisfactory before the peel ply

is removed, it is quite easy to evaluate the surface after removal in time to prevent bonding to an unsatisfactory surface. Figure 7 shows a magnification of the imprint of a nylon peel ply that had supposedly been corona treated to enhance its adhesion to the resin matrix to create a fresh fractured surface over the entire bond. Yet it is quite clear that, at the microscopic level, the surface is absolutely slick. The groove left by every fiber in the peel ply is clearly shown. The matrix is fractured over only an extremely small fraction of the surface, as shown in the higher magnification view of the same surface in Fig. 8. The adhesive will stick only to the fractured surface, not the obviously ultraslick surface so clearly evident in most of Fig. 8. The adhesive could bond to barely 10% of this surface and will not stick anywhere else. A reliable peel ply would leave an imprint so rough at the microscopic level that the weave of the peel ply so prominent in Fig. 7 would never be seen.

FIG. 7 IMPRINT OF A NYLON PEEL PLY LEFT AFTER REMOVAL FROM COMPOSITE LAMINATE. 58×

FIG. 8 IMPRINT OF A PEEL PLY SHOWING EXTREMELY SMALL FRACTION OF SURFACE COVERED BY FRACTURED RESIN MATRIX TO WHICH THE ADHESIVE COULD BOND. 580×

Figure 9 shows the imprint left by a different peel ply, that was believed to be producing satisfactory bonds, or at least much stronger bonds than had been achieved using another peel ply that left a surface like that shown in Fig. 7. There is clearly far less imprint visible in Fig. 9 than in Fig. 7, and far more fractured surface resin. Even so, one would be relying on good adhesion to no more than about 30% of the total area. Also, the surface in Fig. 9 is so rough as to be likely to trap pockets of air between the adhesive and the substrate. Fortunately, unless they coalesced, these air pockets would prevent contact only between the glue and those parts of the laminate to which it would not stick anyway.

FIG. 9 IMPRINT OF A BETTER PEEL PLY ON A COMPOSITE SURFACE SHOWING A FAR GREATER FRACTION OF FRACTURED RESIN OF THE SURFACE. 34×

The real purpose of peel plies is to protect against gross contamination, which is removed as the peel ply is lifted off. Its removal must be followed by grit blasting or by sanding so thorough that all traces of the imprint from the peel ply have been removed. Figure 10 shows a scanning electron micrograph of a carbon-epoxy composite thoroughly sanded with 120 grit abrasive. There is no similarity with either Fig. 7 or Fig. 9. However, if the abrasion is so severe as to expose bare fibers, from which the surface treatment (coating) has been removed everywhere by sanding, the bond will be even

weaker. The solution to this problem is a resin-rich surface layer, like a conventional gel coat applied to simple fiberglass composite components before they are painted.

FIG. 10 SURFACE OF A CARBON-EPOXY COMPOSITE AFTER 120 GRIT ABRASIVE SURFACE PREPARATION. NOTE EXPOSED FIBERS AND SMALL AMOUNT OF DAMAGED FIBERS. 415×

Peel plies are either coated with release agents so that they will not delaminate the composite they are meant to protect or are so slick that the surface is like a mirror at the microscopic level. In either case, the glue cannot stick, except at the small nodules of fractured laminating resin in between the bundles of fibers in the weave of the peel ply. This increment of strength is often sufficient to bond successfully honeycomb cores in sandwich construction. Providing that all the remainder of the processing is done properly, even this limited bond strength is often enough to break the core away from the faces. In addition, this bond strength seems always to be enough to pass any ultrasonic NDI tests, even for fayingsurface bonds. Another reason for thorough abrasion after removal of a peel ply is that the polymer molecules formed adjacent to the peel ply (or released tool surface) are oriented in such a way as to make them very difficult to be wetted by the adhesive. A stronger bond is possible in the more randomly oriented interior molecules. Alternatives to Peel-Ply Surfaces. An alternative technique used to some extent in both Europe and North America is to not use a peel ply as such, but to add an extra layer of preimpregnated fine-weave aramid cloth to both surfaces of the composite laminate, using the same resin as for the primary laminate. This achieves two things. First, it avoids the acutely resin-starved surfaces left as the conventional dry peel plies absorb resin from the base laminate; second, it ensures a completely roughened surface to which to bond, without any trace of release agent from the caul plate.

Tests using fiberglass tear plies that were not intended to separate cleanly, but to remove a little resin from the entire surface, failed because the glass fibers were too brittle and broke repeatedly. The surface was contaminated, and possibly damaged as well, whenever razor blades were used to strip off the surface ply. Obviously, if the tear ply sticks too well, it will not come off alone, but will probably take one or two top plies from the laminate as well. This has been known to occur during the removal of metallic fittings from composite structures that had been installed with a layer of rubber-base sealant between the faying surfaces to prevent galvanic corrosion. Grit Blasting. The objections to grit blasting are easier to recognize, even though they are not all valid. Grit blasting can

create dust--black dust from carbon-epoxy laminates is perceived as a greater health hazard than white dust from fiberglass-polyester laminates. Also, it relies to some extent on operator skill, and if the only available grit-blasting machine must be borrowed from a nearby foundry where it is normally used to clean sand castings, it will be so powerful that it is easy to burn right through a laminate in a matter of seconds. Both of these objections can be overcome, as they were on the Lear Fan all-composite aircraft, by buying an inexpensive low-pressure machine in which the nozzle is surrounded by a ring of vacuum ports, which are themselves surrounded by a circular comb of soft bristles like those on domestic vacuum cleaners. It took only two to three days to properly train someone with no prior mechanical skills to operate such a machine. The bristles held the nozzle a fixed distance away from the work. An adequately blasted surface had a dull matt texture, which contrasted strongly against the shiny black unblasted areas nearby. The pressure was only 140 kPa (20 psi) instead of 700 kPa (100 psi), so it would take such a long time of holding the nozzle in one place to damage the composite surface that it would be obvious that something was wrong long before any damage actually occurred. The grit-blast machine was used in a room that was depressurized slightly with respect to the rest of the factory, but there was no need for protective clothing or respiratory masks. A fixed machine with many nozzles under which the parts pass on a roller table has also been developed. The nozzles may someday be computer controlled. Currently, there are many forms of grit available, some of which are recyclable. Aluminum oxide (Al2O3) has been used traditionally, but is sufficiently expensive that recirculation of the used grit must be considered. It would seem that the use

of recirculated abrasives can be permitted on an otherwise uncontaminated peel-ply surface. However, so much release agent is transferred from caul plates to the adherend surface that no recirculation can be permitted on that kind of surface. Examples of inexpensive media are plastic beads, which can be separated from the debris by floatation in a water tank, biodegradable corn husks, and a number of powders, such as sodium bicarbonate, which can be recovered from aqueous solutions. Since chemical paint stripping is being outlawed, there is considerable pressure to come up with safe replacements. Grit blasting is certainly preferable to the use of jets of hot air, which have been observed to burn as many as two surface plies to the bare fibers in several places on thin-skinned fibrous composite parts. Perhaps the most compelling argument in favor of using grit blasting is that none of the more widely used techniques for initial production can ever be used for repairs, whether by an airline during service or by the original manufacturer because of damage during manufacture. For example, a peel-ply surface cannot be recreated on a previously cured laminate--one is forced to abrade it. Joining of Organic-Matrix Composites L.J. Hart-Smith, McDonnell Douglas Aerospace

The Water-Break Test The water-break test has long been a standard technique for verifying the adequacy of surface preparation for metal bonding. It can also be used for composite surfaces, as described in Ref 4. This is clear from Fig. 11, which compares the even wetting of a properly grit-blasted surface with the repelling of water on an untreated surface contaminated by the release agent on the lay-up tool. Although the water is on the surface for only a small time, and very little is absorbed into the laminate, some adhesives are so moisture-sensitive that it is necessary to dry even that moisture out before bonding. It would seem to make more sense to prime the composite surface after grit blasting so that it could be recleaned using solvents, than to use the water-break test on a surface known to have been prepared properly immediately beforehand. The benefit from the water-break test comes from confirming that a previously prepared panel has not been contaminated during storage.

FIG. 11 COMPARATIVE WATER-BREAK TEST RESULTS. LEFT: PROPER PREPARATION. RIGHT: IMPROPER PREPARATION. COURTESY OF CANADAIR

Experience gained since the preparation of Ref 4 indicates that water-break tests are inappropriate for surfaces that have not been abraded. If water is applied to the surface left by removing a peel ply, the water will bridge between the small areas of fractured resin matrix and totally mask the unwettability of the majority of the surface. The test is ineffective in this case. Worse, it falsely creates the illusion that the surface has been prepared adequately for bonding, as well as the opportunity for someone to shift the blame for a poor bond from failure to dry out water absorbed from other sources to the microscopic traces of water left by the water-break test. Such traces are easily removed by a moderately hot air gun because the water is not on the surface long enough to permit significant absorption.

Matienzo et al. (Ref 13) indicate that lightly sanding the surface created by curing a laminate directly against a tool surface coated with mold release can be good enough to pass the water-break test, yet such surfaces cause bonds weaker than those achieved by untreated peel-ply surfaces. This indicates that even more emphasis must be placed on the thorough sanding or abrasion of the surface to be bonded and less on any subsequent verification. Unfortunately, their test data do not include baseline tests on anodized aluminum adherends against which the composite bonds could be assessed. One of their prime conclusions was that the simple removal of peel plies gave stronger bonds than were achieved with light sanding. However, it is significant that according to their test data, even higher average strengths were attained by deliberately contaminating one of the two peel plies by spraying it with mold release before the panel was cured. It should be pointed out that mold-release agents will stick like glue if they are not baked before a part is cured on the surface. This "contaminated" peel ply gave such better results than the uncontaminated peel ply that it is clear that the untreated peelply surfaces were capable of considerable improvement. Even a peel ply that showed no sign of silicon transfer gave slightly weaker average results than the spray-contaminated peel ply, although the highest individual strength was attained with the silicon-free peel ply. This peel ply was also effective in isolating the laminate from the silicon introduced from the mold release. However, as noted in their work, the use of peel plies on both sides of the part would not be applicable to compound-curved panels or to tool surfaces very much larger than available pieces of peel ply, which could not be lapped without unacceptable mark-off on the outside of the finished part. While even the highest strengths reported in Ref 13 are only about half what one might expect from a grit-blasted surface, the very low strengths they show from either gross contamination or from failure to make any attempt to remove the mold release indicate an irrefutable need for proper surface preparation for bonding polymer composites. The change of surface texture between thoroughly abraded and bare surfaces is so great that, even though the water-break test works reliably, it is not always necessary. A visual inspection, which is capable of distinguishing between thorough and light sanding with the use of a magnifying glass, should suffice. This justification for omitting the water-break test is particularly relevant whenever the adhesive bonding follows immediately after the surface preparation and there is no opportunity for contamination. On the other hand, a drying cycle to remove water applied during the water-break test would also remove absorbed water from near the surface of the part. The omission of a compulsory drying cycle would demand great care in ensuring that the parts to be bonded had not absorbed water since their manufacture, even if they had been protected against the more widely recognized contaminants. The water-break test would still seem to be very valuable for repairs to structures contaminated by chemicals absorbed during service or by coupling fluid for ultrasonic NDI of a porous or delaminated panel. It has already been shown to be effective in proving that most cleaning solvents applied to an unprimed surface actually contaminate the surface and need to be removed by water-soluble solvents such as alcohol. This issue is discussed in Ref 4, which includes photographs of the beach marks left as the solvent evaporates. Experience with using the water-break test has revealed that the application of the least amount of water possible is preferable because a thick film of water can bridge a small local area of contamination and conceal a defect such as a thumb print. One needs to watch the film as it evaporates. A slow uniform drying out simultaneously over the entire surface is an indication of satisfactory surface preparation. If, on the other hand, a discrete dry area suddenly appears at some point during the evaporation, that is an indication of contamination.


4. L.J. HART-SMITH, R.W. OCHSNER, AND R.L. RADECKY, SURFACE PREPARATION OF FIBROUS COMPOSITES FOR ADHESIVE BONDING OR PAINTING, DOUGLAS SERVICE MAGAZINE, FIRST QUARTER 1984, P 12-22 13. L.J. MATIENZO, J.D. VENABLES, J.D. FUDGE, AND J.J. VELTEN, SURFACE PREPARATION OF BONDING ADVANCED COMPOSITES, PART 1; EFFECT OF PEEL PLY MATERIALS AND MOLD RELEASE AGENTS ON BOND STRENGTH, 30TH NATIONAL SAMPE SYMPOSIUM, SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 19-21 MARCH 1985, P 302314

Joining of Organic-Matrix Composites L.J. Hart-Smith, McDonnell Douglas Aerospace

Control of the Bonding Process for Thermoset Adhesives The chemistry of creating a properly cross-linked adhesive bond is well understood by a large number of materials and process engineers. However, far fewer people who actually control autoclaves are knowledgeable of adhesive chemistry or have real-time access to such information. With the help of a polymer chemistry expert, an experienced process engineer can identify the cause of a problem and modify a cure cycle in progress to save a part that would otherwise need to be scrapped. Some of the more important process variables that need to be understood and controlled are outlined below. Effects of Temperature Variations on Bond Strength. Some aspects of the process specification are almost so self-

evident that their significance is underestimated. It is obvious that if an adhesive is heated up too slowly, it will not wet the surface to be bonded. The adhesive may eventually cure, but there will be a dry bond. Therefore, the process specifications include minimum heat-up rates. These vary with the adhesive, but are on the order of 3 °C (5 °F) per minute. Likewise, an adhesive can be made to boil if heated up very much faster, so there must be maximum heat-up rates and maximum temperatures as well. These limits present no problem when making a small, uniformly thick test panel. The problems arise when a large, complex assembly of variable thickness is either cured or bonded on a large tool having a variable heat sink. It then becomes very easy for some parts of the structure to be heated up too slowly while other parts are on the verge of being burned. For all but the simplest of parts, a proof-of-tool demonstrator should be carried out to ensure that the assembly can be properly cured the first time. It makes sense to save on the cost of this operation by using cheaper fibers with the same resin. One solution for the problem of a large range in temperature within the composite part or assembly, once a cure is already in progress, is to put constant-temperature dwells in the cure cycle to give the various segments of the part time to equilibrate. This is best done early in the cure cycle, before any of the adhesive (or resin matrix) has gelled. Thus, in addition to maximum and minimum heat-up rates, there is a need for a maximum variation between high and low temperatures--measured on the part, not on the tool, even if that takes more time than it does to hook up a permanently wired thermocouple. In some cases, multiple dwells may be needed. There is a limit, however, since an excessive number of dwells will cause the adhesive to gel before it has a chance to wet the surface needing to be bonded. Once such a problem is detected, blankets can be added for subsequently made parts to shield the hot spots and local heaters added to overcome the colder areas. It is not acceptable to simply disconnect the thermocouples giving the highest and lowest readings in subsequent cure cycles. It is better still to thermally isolate large heat sinks from the rest of the tool if this can be done without causing distortion of the part. A simple but effective solution to some such problems is to thermally isolate the tool from the bed of the autoclave and, wherever possible, to thermally isolate the tool surface contacting the part or assembly from the supporting structure. Other than the possibility of tool distortion, which may or may not adversely affect the product, no harm will come to the composite part or bonded assembly if it is heated correctly and only the tool support structure lags behind. Most tools are made from metal because hand layed-up plastic tools tend to have internal voids that collapse and need patching. However, the metal tools expand and contract more than composite parts. A solution to this problem is to use autoclave-cured composite members to hold composite parts in the work station during a bond cycle. The composite holding frame is supported by the basic metal tool, but it is held at only one point and permitted to slide freely as the size of the metal tool changes. This enables linear dimensions to be held and avoids the problem of crushing the composite parts as the tool contracts during cool down. It also thermally isolates composite parts from the metal tools. The problem of tools having excessive heat sinks arises because tool designers are used to extremely rigid tools rather than tools that are barely able to maintain their shape, but so simple and compact as to permit many of them to be loaded into a single autoclave cycle. For thin, flexible skins, it makes more sense to rack several caul plates in a single holding fixture than to make one massive lay-up and cure tool per part. The productivity of the autoclave is enhanced this way as well. While the caul plates may distort out-of-plane, they will retain their in-plane dimensions, and only light pressure will be needed to restore the profile of the parts as they are bonded into assemblies. For example, 9 m (30 ft) long Lear Fan unstiffened fuselage skin panels, which were mildly compound curved, were rolled up like giant cigars to stiffen

them for safer as well as easier handling as they were transported from the autoclave to the assembly room. The more common flat panels flex even more easily. Obviously, the entire shape must be well defined for the final bonding operation since the assembly will be too rigid to match the surrounding structure if it is made to the wrong shape. However, the bond cure cycle is typically 25 to 40 °C (50 to 100 °F) lower in peak temperature than the cure cycle for detail composite parts, so the problems of variable temperatures are diminished. Effects of Prebond Moisture on the Strength of Bonded Joints. Another widespread problem is the need for the

removal of volatiles before the start of cure and throughout its duration. Some manufacturers of composite and bonded structures want materials that are so tolerant of abusive cure cycles that it is permissible to: (1) apply full pressure from the start of the cure, flipping one switch for the heaters (with no automatic control system); and (2) to come back once during the cure at a preestablished time to cut back on the heat when the peak temperature has been attained and once more to switch the heat off when the cure is complete, preferably quickly enough not to need the services of a second shift of workers. Material development engineers have failed to create such forgiving materials, even though some process specifications permit parts to be made this way, so there are opportunities for using closed-loop control systems. Such systems, however, require experts to anticipate what should happen, what might happen, and what must not be permitted to happen--along with corrective responses. The reliability of manufacture will be increased by demanding that all materials and parts be dried thoroughly immediately before each cure cycle. This refers also to Nomex honeycomb cores, which are known to have a great affinity for water and to absorb it rapidly if left in a humid environment. Wet cores result in weak glue fillets at honeycomb cell walls (Fig. 12). Such a goal of thorough drying exceeds the current practice at many composite manufacturing facilities today, whether or not the up-front savings exceed the consequently higher down-stream costs. But even if thorough drying were the norm, bagging procedures that permitted small amounts of moisture to be removed harmlessly during cure would also be desirable. The initial manufacturer also needs to be aware of the very high costs (to the consumer and/or the manufacturer) whenever a composite part or assembly is manufactured with sufficient trapped moisture to create an inferior part of unreliable service strength, but sufficient initial strength to be accepted after normal ultrasonic inspection. Under the prevailing circumstances, the original manufacturer is exposed to the consequences of improper processing only when there is sufficient residual strength in a rejected bonded panel to make it virtually impossible to disassemble the pieces for rebonding without destroying them by delaminations in the process.

FIG. 12 EFFECT OF NOT PREDRYING ORGANIC HONEYCOMB CORES. (A) DRIED CORE. (B) UNDRIED CORE

Adhesives vary in their susceptibility to problems caused by moisture absorption. The first generation of modified epoxy adhesives introduced in the 1960s were discredited because of corrosion problems with bonded metal structures. These adhesives had a high permeability for moisture absorption after cure. This led to corrosion of the metal surfaces, which were usually only etched, not anodized, and were not protected by a corrosion-inhibiting primer. Ironically, these same adhesives have been used successfully with composite adherends that cannot corrode, but they have suffered limited application because of their bad reputation earned in a totally unrelated context. Actually, it was proved more than 20 years ago that the primary cause of such problems was inadequate surface preparation. However, other contemporary adhesives, cured at higher temperatures, lasted longer on the same poor surfaces and were then used in their place.

The second generation of modified epoxy adhesives, introduced in the late 1970s, were specially formulated to impede moisture migration and were used in combination with a phenolic-based primer baked on an aluminum alloy surface anodized with phosphoric acid. Their durability was established during the PABST program, and they have performed well in metal bonded structures. However, this class of adhesives has a great affinity for moisture in the uncured state. One of the first developed was so sensitive that it was withdrawn from sale. Other adhesives in this class are now associated with numerous problems with bonding composite laminates together and with the co-cure and bonding of composite patches. The problem, which was exposed and evaluated during the PABST program, occurs only when the adhesive is allowed to absorb moisture before bonding or when making sandwich panels with wet foam or Nomex honeycomb cores. The adhesive film can be restored by drying before use, but the best solution is to avoid the problem by limiting the exposure to humid environments or by regulating the atmosphere in the lay-up room. It is vital that adhesives be stored in relatively small rolls so that the supply of material for an excessive lay-up period is not exposed to the atmosphere. They should be removed from frozen storage the day before they are to be used and allowed to thaw out overnight while still sealed, to prevent condensate from forming on the adhesive. If they do absorb moisture, they should be dried before further use. It is better to close down production for a day than to take a roll of adhesive straight out of the freezer, unwrap it, and start using it immediately while it is covered with condensate. Laboratory experiments performed to prove that trapped moisture could actually prevent the adhesive from sticking, not merely degrade its properties, required fairly large blister panels to replicate the effect. It was shown that 150 mm (6 in.) square panels that were sealed around the edges would show the effect, but that smaller ones would not. It follows, therefore, that the conventional half-square-inch lap-shear test coupons used as travellers while parts are made cannot possibly expose this condition. If the adhesive has been exposed to high humidity and temperature for even the normal duration of a lay-up operation, or covered with condensate because it was unsealed before it had thawed out, the weak bonds will almost certainly be in the structure. If left out for an entire week, them is a good chance that the pieces will fall apart as they are removed from the bonding tool. What is far worse than the above scenario is a degraded adhesive bond with sufficient residual strength that it does fall apart. There is then no way of characterizing the strength or the time taken for the bonds to fail in service. They will pass all ultrasonic inspections, unless there was so much trapped water that large blisters occurred during cure. Better composite parts can be made more cheaply if the manufacturer starts with a slow vacuum-consolidation cycle, usually overnight, not in an autoclave, but possibly in a low-temperature oven or under an insulating styrofoam shelter. This enables trapped air bubbles to escape, solvents to be driven off, and some moisture to be expelled. Failure to do this, or the equivalent step of thoroughly drying detail parts to be bonded, as well as the adhesive itself if it is suspected of absorbing moisture, guarantees that all of these ingredients will be trapped inside the part once the autoclave pressure is applied. Full vacuum would be applied inside the bag until slightly short of the minimum viscosity point for the resin matrix, or until no further volatiles could be detected in the vent, if that occurred sooner. Most resins will boil if cured under high vacuum, so the bag needs to be vented to the atmosphere once the vent is no longer useful. The cure can then proceed, with the autoclave gage pressure holding the laminate down and the resin in the bag being subjected to the atmospheric pressure. After gelation, the pressure felt by the rosin will increase progressively toward that experienced inside the autoclave. Partial-Vacuum Cures. A minor variation of the scheme above can eliminate the need for autoclaves, when curing

simple parts that need little pressure to hold the laminates against the tool surface. After an unusually thorough consolidation cycle, the cure would start with a full vacuum inside the bag, using only an oven as a heat source. Once the point of minimum viscosity was reached, a regulated pressure of between one-third and two-thirds of an atmosphere would be maintained inside the bag, with the remaining differential pressure holding the laminate onto the tool. Many, but not all, resins should be capable of being cured this way. Phenolics, which generate steam during cure, are the most obvious exception. There are new adhesives and resins being developed that will tolerate curing under full vacuum throughout the entire cycle, but the approach outlined above would produce better laminates and bonds even with such materials. It would also eliminate the need for an autoclave for many more resins and adhesives than a requirement for full vacuum would permit. It is a technique worthy of further study, particularly for repairs to thin composite sections. In-house assessments of both new and old adhesives at Douglas Aircraft have shown that the partial-vacuum cure achieves bond strengths almost as great as those produced by full autoclave pressure, and far better than those achieved under a full-vacuum cure of the same adhesive. This is true of adhesives specially developed for vacuum-bag curing as well as those formulated for autoclave cures and known to be totally unsatisfactory if cured under full vacuum. It is likely that the need for more heat

to maintain the heat-up rate on thicker panels, or with large tools that have not been thermally isolated from the work, would favor the use of an autoclave because the thermal conductivity of air increases as it is compressed. Susman (Ref 14) has described the benefits of the dwell cycle for autoclave cures. Maintaining full vacuum while holding the temperature for about an hour at some 55 °C (100 °F) below the cure temperature permitted such a thorough removal of volatiles that the interlaminar shear strengths were doubled with respect to traditional cures. The worst laminates were made by the so-called straight-up cure, in which both full heat and full autoclave pressure were applied at the start of cure to minimize the cycle time.


14. S.E. SUSMAN, THE DWELL CYCLE: A COST SAVINGS ROUTE TO QUALITY COMPOSITES, 25TH NATIONAL SAMPE SYMPOSIUM, SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 6-8 MAY 1980, P 251-258 Joining of Organic-Matrix Composites L.J. Hart-Smith, McDonnell Douglas Aerospace

Bonding of Thermoplastic Composite Panels Three different approaches have been evaluated for bonding thermoplastic composite components together. Adhesive bonding has been used, much as for thermoset composites, but needing different surface preparations. Fusing the panels together with or without an interlayer has also been used. This technique is of interest to engineers because it suggests that repairs could be performed far more easily than is possible with thermoset composites. However, there is the obvious disadvantage that it is difficult to melt such a localized area of resin without distorting the rest of the structure. Perhaps the most promising approach is that of coating the surfaces to be bonded with a different thermoplastic, which melts at a lower temperature than that in the basic laminate. Fusion bonding is then possible at a temperature low enough not to distort the part itself. Bonding with Thermoset Adhesives. The reluctance of liquid polymers to wet solid thermoplastics is even greater than for thermosets. Having the surfaces free from contamination is but one of many criteria that must be met for one material to stick to another. Having the right surface energies is equally critical. An uncured resin or adhesive will typically not wet a cured surface having a far lower surface tension, as is achieved with release coatings. It is much more likely to adhere to a surface of far higher surface energy, as can be created by oxidation via corona treatment. Figure 13 shows how little corona energy is needed to almost double the surface energy of polyether ether ketone (PEEK). However, testing also showed that far more corona energy was needed to ensure a cohesive rather than interfacial failure (Ref 8). This supports the assessment that even after surface contaminants have been removed, there are still polymer layers of low surface energy between the bulk matrix and the surface created by curing against an inert tool surface. Much the same problem exists when trying to epoxy bond injection-molded thermoplastic inserts into honey-comb sandwich panels. The glue will not stick unless the initially inert plastic surface has been activated to form a chemical bond or etched to permit a mechanical bond.

FIG. 13 EFFECT OF CORONA TREATMENT ON THE SURFACE ENERGY OF PEEK. SOURCE: REF 8

Wu et al. (Ref 15) describe the adhesive bonding of thermoplastic composites using several adhesives and a variety of surface treatments. Figure 14 shows results from their study. It is clear from Fig. 14 that weak interfacial failures are the inevitable result of failure to appreciate the need for proper surface treatment, for both the hot-bond epoxy FM-300 film adhesive and the BMI film adhesive EA-9673. Cohesive failures within the bond layer were achieved only for the chromic-sulfuric acid etch and plasma treatments. Significantly, these are the highest strengths obtained. While the highest average shear strengths are slightly less than one would expect on anodized aluminum coupons, they are credible for fibrous composite adherends having a quasi-isotropic fiber pattern and a slightly lower modulus. Higher strengths should have been obtained on unidirectional test coupons.

FIG. 14 EFFECT OF SURFACE TREATMENT ON THE LAP SHEAR STRENGTH OF AS4-APC-2 LAMINATE. SOURCE:

REF 15

Other researchers have also reported success with plasma treatment, with various gaseous environments. While the technique is effective, and probably quite practical for small parts, the need to enclose the composite part in an evacuated chamber during the surface preparation would seem to render it impractical for very large components. On the other hand, the reliability of this process in comparison with others is of paramount concern. The acid etching would require a large tank, in any event, and the corona-discharge technique would require electrodes tailored for each part to achieve uniform treatment over the area if the thickness varied or if the part were contoured. The capital cost of the equipment for carrying out the process correctly is a small fraction of the cost of scrapping one or two large composite assemblies because the surfaces had not been prepared properly while there was access to the individual panels. Figure 15 makes it clear that mere abrasive cleaning of thermoplastic adherends is unsatisfactory, no matter what thermoset adhesive is used, while the chromic-sulfuric acid etch is effective with all the adhesives tested. In other tests, Wu and co-workers found that extended etching did not increase the joint strength; worse, it pitted and degraded the surface. They also established by preliminary testing that them is a need to consider the environmental durability of bonds to thermoplastic adherends, just as them is with thermoset adherends.

FIG. 15 EFFECT OF ABRASIVE CLEANING AND ACID ETCHING ON THE LAP SHEAR STRENGTH OF AS4/APC-2 LAMINATES. SOURCE: REF 15

Fusion Bonding of Thermoplastic Composites. Wu and co-workers also included some welded thermoplastic joints

in their study (Ref 15). Lap-shear coupons that were bonded by ultrasonic welding exhibited good strengths for this process, whether the surface had been acid etched or merely abraded (Fig. 16). However, the procedure was not considered practical for large structures. Wu and co-workers examined such techniques as resistance welding, induction welding, laser welding, hot plate welding, and ultrasonic welding. It was determined that the high carbon-fiber content in the thermoplastic matrix and the high thermal and electrical conductivities of the fibers impose a tremendous amount of difficulty on fusion bonding processing. Uneven heating throughout the bonded area and delamination and distortion of the composite materials must be overcome with fusion bonding methods. In addition, such welding methods will change the crystalline structure of the resin because the remelting would be followed by cooling at a different rate from that experienced when the laminate was originally made.

FIG. 16 EFFECT OF ABRASIVE CLEANING AND ACID ETCHING ON FUSION-BONDED AS4-APC-2 LAMINATES. TEST COUPONS WERE JOINED BY ULTRASONIC WELDING. SOURCE: REF 15

Fusion Bonding Using the Dual Resin Approach. The problems cited above can be avoided by using an interlayer

that melts at a lower temperature than the parent thermoplastic matrix, but which is miscible with it. Several studies have been made using polyetherimide (PEI) to join PEEK laminates. Zelenak and Radford (Ref 16) cite success with a process in which a PEI film was melted onto the surface of APC2/PEEK tubes prior to joining. The PEI flowed into the PEEK parent, becoming a part of the parent material. This process occurred at the 320 °C (610 °F) processing temperature for the PEEK. To form the actual joint, another film of PEI was placed between the pieces to be joined as a filler. The joint was made under light pressure at a temperature of 300 °C (570 °F). This temperature is substantially above the glass transition temperature of PEI, which is 210 °C (410 °F), but below the melting temperature of PEEK (334 °C, or 633 °F). At this temperature the joint was formed by the PEI film bonding with the PEI in the parent material, which was intimately mixed with the PEEK material. Unfortunately, the testing was marred by the limited strength of the structure outside the joint which, while it may have been able to carry the design loads, was not strong enough to fail either the laminate adjacent to the joint or the bond itself, except for two premature failures. Consequently, the absolute strengths of the joints were not established. However, in other regards, the use of PEI/PEEK as a miscible combination for bonding of structural PEEK matrix composites gives excellent joint performance with minimal susceptibility to surface contamination during bonding. While there remains the question of resistance to the environment, various chemicals, and age, this method is most promising. It would obviously be feasible to use the same technique to either repair damage in service or to add additional bonds as requirements changed, provided that the entire surface of the parts had been coated with PEI when the details were initially made. Yoon and McGrath (Ref 17) have reported on bonding PEEK laminates with polyimide thermoplastic adhesive films. Their work is similar to that above in terms of processing at temperatures insufficient to melt the PEEK, but it differed significantly in that they bonded directly to the PEEK laminate. The absence of a compatible surface layer caused them to investigate the effectiveness of different surface treatments. They were able to make very good bonds without adhesive at 360 °C (680 °F), developing shear strengths of 47 MPa (6800 psi), which are not surprising given that their adherends were 24-ply thick unidirectional carbon-fiber reinforced laminates. However, they concluded that this temperature could not be used due to severe deformation. Their joints that fused together at 340 and 350 °C (645 and 660 °F) attained less than 20% of the highest strengths reached with adhesive films fused to treated surfaces that had been promptly primed with a solution of the adhesive to protect the surface. It was concluded that the highest practical temperature for the PEEK-graphite composite adhesion was around 340 °C (645 °F); however, high pressure was also required. Simple fusion of thermoplastic laminates does not seem to be a productive avenue at the present time. It should be pointed out, however, that the work of Yoon and McGrath in activating the surface prior to bonding with thermoplastic adhesive closely parallels what was reported earlier for bonding with thermoset adhesives. Grit blasting showed a distinct improvement but clearly failed to enable the true material strengths to be exploited. The various gas plasma treatments were effective in producing cohesive failures, and the combinations of grit blasting and gas plasma treatments were better still, reaching nearly 41 MPa (6000 psi) and coming close to the strengths attained by fusion at 360 °C (680 °F). Of the successful techniques described above, only one is applicable for joining large components. Fusion welding at a reduced temperature, made possible because the thermoplastic parts have been coated during fabrication with a miscible

plastic that melts at a lower temperature than does the base matrix, should enable full use to be made of the strength of the parent laminates. The plasma treatment and corona-discharge techniques are reliable for small parts. The other techniques for developing high strengths are nevertheless invaluable as a warning of how weak the structures will be if no attention is paid to the need for proper surface preparation. They establish the target that should be obtained by any method used in practice. Solvent Bonding of Amorphous Thermoplastics. The difficulties cited above, when bonding aircraft-grade fiberreinforced thermoplastic composites, apply specifically to crystalline thermoplastics. Amorphous thermoplastics have not been used widely for such structures because of their lower resistance to common solvents. However, this greater susceptibility to solvents can become an advantage for adhesive bonding. The extreme case of assembling polystyrene plastic model aircraft kits by local application of styrene monomer cement or methyl ethyl ketone solvent is well known. Solvent bonding of nonstructural thermoplastics is quite common in other industries. The same effect can also be used to enhance the bonding of amorphous thermoplastics by epoxy-type adhesives. Such adhesives do not normally bond well to these plastics. However, wiping the surfaces with a suitable solvent and bonding the parts together before the thermoplastic has time to revert to an inert state can produce adhesive bonds strong enough to tear honeycomb cores apart; that is, strength similar to good metal-bond structures. This has been demonstrated by the Ten Cate organization for their continuous-fiber-reinforced PEI thermoplastic bonded with 120 °C (250 °F) cured modified epoxy adhesives. They even advocate the direct use of solvent bonding for cosmetic repairs of local damage to such structures. This technique is no panacea, however, because some thermoplastics will craze after being subjected to the wrong solvent. Nevertheless, under suitable circumstances, solvent bonding or solvent-enhanced bonding can be a very useful technique.


8. A.J. KINLOCH AND G.K.A. KODOKIAN, "THE ADHESIVE BONDING OF THERMOPLASTIC COMPOSITES," EUROPEAN RESEARCH OFFICE OF THE U.S. ARMY, CONTRACT DAJA 45-86-C0037, FOURTH PERIODIC REPORT, IMPERIAL COLLEGE, LONDON, AUG-DEC, 1987 15. S-I.Y. WU, A.M. SCHULER, AND D.V. KEANE, ADHESIVE BONDING OF THERMOPLASTIC COMPOSITES, 1. THE EFFECT OF SURFACE TREATMENT ON ADHESIVE BONDING, 19TH INT. SAMPE TECH. CONF., SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 13-15 OCT 1987, P 277-290 16. S. ZELENAK, D.W. RADFORD AND M.W. DEAN, THE PERFORMANCE OF CARBON FIBER REINFORCED PEEK SUBASSEMBLIES JOINED USING A DUAL RESIN BONDING APPROACH, 37TH INTERNATIONAL SAMPE SYMPOSIUM, SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 9-12 MARCH 1992, P 1346-1356 17. T-H. YOON AND J.E. MCGRATH, ADHESION STUDY OF PEEK/GRAPHITE COMPOSITES, 36TH INTERNATIONAL SAMPE SYMPOSIUM, SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 15-18 APRIL 1991, P 428-436 Joining of Organic-Matrix Composites L.J. Hart-Smith, McDonnell Douglas Aerospace

References

1. L.J. HART-SMITH, JOINTS, COMPOSITES, VOL 1, ENGINEERED MATERIALS HANDBOOK, T.J. REINHART, ED., ASM INTERNATIONAL, 1987, P 479-495 2. E.W. THRALL, JR. AND R.W. SHANNON, ED., ADHESIVE BONDING OF ALUMINUM ALLOYS, MARCEL DEKKER, 1985 3. R.J. SCHLIEKELMANN, ADHESIVE BONDING AND COMPOSITES, PROGRESS IN SCIENCE AND ENGINEERING OF COMPOSITES, VOL 1, PROC. 4TH INT. CONF. COMPOSITE MATERIALS (TOKYO), 22-28 OCT 1982, JAPAN SOCIETY FOR COMPOSITE MATERIALS AND THE

METALLURGICAL SOCIETY (TMS) OF AIME, T. HAYASHI, K. KAWATA, AND S. UMEKAWA, ED., NORTH-HOLLAND, AMSTERDAM, 1982, P 53-78 4. L.J. HART-SMITH, R.W. OCHSNER, AND R.L. RADECKY, SURFACE PREPARATION OF FIBROUS COMPOSITES FOR ADHESIVE BONDING OR PAINTING, DOUGLAS SERVICE MAGAZINE, FIRST QUARTER 1984, P 12-22 5. L.J. HART-SMITH, R.W. OCHSNER, AND R.L. RADECKY, SURFACE PREPARATION OF FIBROUS COMPOSITES FOR ADHESIVE BONDING OR PAINTING, CANADAIR SERVICE NEWS, VOL 14 (NO. 2), SUMMER 1985, P 2-8 6. B.M. PARKER AND R.M. WAGHORNE, SURFACE PRETREATMENT OF CARBON FIBREREINFORCED COMPOSITES FOR ADHESIVE BONDING, COMPOSITES, VOL 13, JULY 1982, P 280-288 7. A.V. POCIUS AND R.P. WENZ, MECHANICAL SURFACE PREPARATION OF GRAPHITE-EPOXY COMPOSITE FOR ADHESIVE BONDING, SAMPE J., SEPT/OCT 1985, P 50-58 8. A.J. KINLOCH AND G.K.A. KODOKIAN, "THE ADHESIVE BONDING OF THERMOPLASTIC COMPOSITES," EUROPEAN RESEARCH OFFICE OF THE U.S. ARMY, CONTRACT DAJA 45-86-C0037, FOURTH PERIODIC REPORT, IMPERIAL COLLEGE, LONDON, AUG-DEC, 1987 9. J.D. LABOR AND S.H. MYHRE, "LARGE AREA COMPOSITE STRUCTURE REPAIR," NORTHROP AIRCRAFT CORPORATION, USAF TECHNICAL REPORT AFFDL-TR-79-3040, MARCH 1979 10. B.M. PARKER, THE EFFECT OF COMPOSITE PREBOND MOISTURE ON ADHESIVE-BONDED CFRP-CFRP JOINTS, COMPOSITES, VOL 14, JULY 1983, P 226-232 11. E.W. THRALL, JR., PROSPECTS FOR BONDING PRIMARY AIRCRAFT STRUCTURES IN THE 80'S, 25TH NATIONAL SAMPE SYMPOSIUM, VOL 25, SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 6-8 MAY 1980, P 716-727 12. I. MCARTHUR, HAWKER DEHAVILLAND, BANKSTOWN, AUSTRALIA, PRIVATE COMMUNICATION, 1992 13. L.J. MATIENZO, J.D. VENABLES, J.D. FUDGE, AND J.J. VELTEN, SURFACE PREPARATION OF BONDING ADVANCED COMPOSITES, PART 1; EFFECT OF PEEL PLY MATERIALS AND MOLD RELEASE AGENTS ON BOND STRENGTH, 30TH NATIONAL SAMPE SYMPOSIUM, SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 19-21 MARCH 1985, P 302314 14. S.E. SUSMAN, THE DWELL CYCLE: A COST SAVINGS ROUTE TO QUALITY COMPOSITES, 25TH NATIONAL SAMPE SYMPOSIUM, SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 6-8 MAY 1980, P 251-258 15. S-I.Y. WU, A.M. SCHULER, AND D.V. KEANE, ADHESIVE BONDING OF THERMOPLASTIC COMPOSITES, 1. THE EFFECT OF SURFACE TREATMENT ON ADHESIVE BONDING, 19TH INT. SAMPE TECH. CONF., SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 13-15 OCT 1987, P 277-290 16. S. ZELENAK, D.W. RADFORD AND M.W. DEAN, THE PERFORMANCE OF CARBON FIBER REINFORCED PEEK SUBASSEMBLIES JOINED USING A DUAL RESIN BONDING APPROACH, 37TH INTERNATIONAL SAMPE SYMPOSIUM, SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 9-12 MARCH 1992, P 1346-1356 17. T-H. YOON AND J.E. MCGRATH, ADHESION STUDY OF PEEK/GRAPHITE COMPOSITES, 36TH INTERNATIONAL SAMPE SYMPOSIUM, SOCIETY FOR THE ADVANCEMENT OF MATERIAL AND PROCESS ENGINEERING, 15-18 APRIL 1991, P 428-436 Joining of Oxide-Dispersion-Strengthened Materials David O'Donnell, Inco Alloys International Inc.

Introduction OXIDE - DISPERSION - STRENGTHENED (ODS) MATERIALS utilize an extremely fine oxide dispersion for strengthening, such as yttria (Y2O3) or alumina (Al2O3). The majority of ODS alloys are produced by mechanical means, often referred to as mechanical alloying (MA), where elemental or alloyed powders are ball milled, or attrited. This mechanical working of different powders between large shot will cold weld particles and break them apart until a fine interlayered structure is produced. The powder is typically consolidated and homogenized by hot isostatic pressing or hot extrusion. Further consolidation is typically performed, producing a final product that is more than 99% dense (Ref 1, 2). The compositions of selected ODS alloys are given in Table 1.

TABLE 1 COMPOSITION OF SELECTED ODS ALLOYS TYPICAL APPLICATIONS

ODS ALLOY MA 956

COMPOSITIONS, % Cr Ni Fe Ti 20.0 . . . 74.0 0.5

Al 4.5

C ...

Y2O3 0.5

OTHER ...

PM 2000 MA 754

20.0 20.0

... 78.0

73.0 1.0

0.5 0.5

5.5 0.3

... 0.05

0.5 0.6

... ...

PM 1000 MA 758 MA 6000

20.0 30.0 15.0

79.0 68.0 69.0

... 1.0 ...

0.5 0.5 2.5

0.3 0.3 4.5

... ... 0.05

0.6 0.6 1.1

PM 3030

17.0

68.2

...

...

6.0

25 mm, or 1.0 in.) composite compression-loaded submersible is a block compression joint (Fig. 14). The type of loading for a submersible structure is primarily hydrostatic pressure, which results in biaxial compression loading of 2:1 for cylinders and equal biaxial loading for spheres. Bolted joints are not recommended for submersible structures because they provide a path for a potential water leak. Submersibles made from composites for significant depth operation require thick composite walls that cannot be readily joined by adhesive bonding. Bonded joints are best suited for joining thin structures (

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