Properties and Selection: Irons, Steels, and High Performance Alloys

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01 Sep 2005

Copyright ASM International. All Rights Reserved.

Page 1

ASM Handbook,Volume 1

Publication Information and Contributors

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Publication Information and Contributors Publication Information and Contributors Properties and Selection: Irons, Steels, and High-Performance Alloys was published in 1990 as Volume 1 of the 10th Edition Metals Handbook. With the second printing (1993), the series title was changed to ASM Handbook. The Volume was prepared under the direction of the ASM International Handbook Committee.

Authors and Reviewers G. Aggen Allegheny Ludlum Steel Division Allegheny Ludlum Corporation Frank W. Akstens Industrial Fasteners Institute C. Michael Allen Adjelian Allen Rubeli Ltd. H.S. Avery Consultant P. Babu Caterpillar, Inc. Alan M. Bayer Teledyne Vasco Felix Bello The WEFA Group S.P. Bhat Inland Steel Company M. Blair Steel Founders' Society of America Bruce Boardman Deere and Company Technical Center Kurt W. Boehm Nucor Steel Francis W. Boulger Battelle-Columbus Laboratories (retired) Greg K. Bouse Howmet Corporation John L. Bowles North American Wire Products Corporation J.D. Boyd Metallurgical Engineering Department Queen's University B.L. Bramfitt Bethlehem Steel Corporation Richard W. Bratt Consultant W.D. Brentnall Solar Turbines


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01 Sep 2005

C.R. Brinkman Oak Ridge National Laboratory Edward J. Bueche USS/Kobe Steel Company Harold Burrier, Jr. The Timken Company Anthony Cammarata Mineral Commodities Division U.S. Bureau of Mines A.P. Cantwell LTV Steel Company M. Carlucci Lorlea Steels Harry Charalambu Carr & Donald Associates Joseph B. Conway Mar-Test Inc. W. Couts Wyman-Gordon Company Wil Danesi Garrett Processing Division Allied-Signal Aerospace Company John W. Davis McDonnell Douglas R.J. Dawson Deloro Stellite, Inc. Terry A. DeBold Carpenter Technology Corporation James Dimitrious Pfauter-Maag Cutting Tools Douglas V. Doanne Consulting Metallurgist Mehmet Doner Allison Gas Turbine Division Henry Dormitzer Wyman-Gordon Company Allan B. Dove Consultant (deceased) Don P.J. Duchesne Adjelian Allen Rubeli Ltd. Gary L. Erickson Cannon-Muskegon Corporation Walter Facer American Spring Wire Company Brownell N. Ferry LTV Steel Company F.B. Fletcher Lukens Steel Company E.M. Foley Deloro Stellite, Inc. R.D. Forrest Division Fonderie Pechinery Electrometallurgie James Fox Charter Rolling Division Charter Manufacturing Company, Inc.


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01 Sep 2005

Edwin F. Frederick Bar, Rod and Wire Division Bethlehem Steel Corporation James Gialamas USS/Kobe Steel Company Jeffery C. Gibeling University of California at Davis Wayne Gismondi Union Drawn Steel Co., Ltd. R.J. Glodowski Armco, Inc. Loren Godfrey Associated Spring Barnes Group, Inc. Alan T. Gorton Atlantic Steel Company W.G. Granzow Research & Technology Armco, Inc. David Gray Teledyne CAE Malcolm Gray Microalloying International, Inc. Richard B. Gundlach Climax Research Services I. Gupta Inland Steel Company R.I.L. Guthrie McGill Metals Processing Center McGill University P.C. Hagopian Stelco Fastener and Forging Company J.M. Hambright Inland Bar and Structural Division Inland Steel Company K. Harris Cannon-Muskegon Corporation Hans J. Heine Foundry Management & Technology W.E. Heitmann Inland Steel Company T.A. Heuss LTV Steel Bar Division LTV Steel Company Thomas Hill Speedsteel of New Jersey, Inc. M. Hoetzl Surface Combustion, Inc. Peter B. Hopper Milford Products Corporation J.P. Hrusovsky The Timken Company David Hudok Weirton Steel Corporation S. Ibarra Amoco Corporation J.E. Indacochea Department of Civil Engineering, Mechanics, and Metallurgy


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University of Illinois at Chicago Asjad Jalil The Morgan Construction Company William J. Jarae Georgetown Steel Corporation Lyle R. Jenkins Ductile Iron Society J.J. Jonas McGill Metals Processing Center McGill University Robert S. Kaplan U.S. Bureau of Mines Donald M. Keane LaSalle Steel Company William S. Kirk U.S. Bureau of Mines S.A. Kish LTV Steel Company R.L. Klueh Metals and Ceramics Division Oak Ridge National Laboratory G.J.W. Kor The Timken Company Charles Kortovich PCC Airfoils George Krauss Advanced Steel Processing and Products Research Center Colorado School of Mines Eugene R. Kuch Gardner Denver Division J.A. Laverick The Timken Company M.J. Leap The Timken Company P.W. Lee The Timken Company B.F. Leighton Canadian Drawn Steel Company R.W. Leonard USX Corporation R.G. Lessard Stelpipe Stelco, Inc. S. Liu Center for Welding and Joining Research Colorado School of Mines Carl R. Loper, Jr. Materials Science & Engineering Department University of Wisconsin-Madison Donald G. Lordo Townsend Engineered Products R.A. Lula Consultant W.C. Mack Babcock & Wilcox Division McDermott Company T.P. Madvad USS/Kobe Steel Company


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J.K. Mahaney, Jr. LTV Steel Company C.W. Marshall Battelle Memorial Institute G.T. Matthews The Timken Company Gernant E. Maurer Special Metals Corporation Joseph McAuliffe Lake Erie Screw Corporation Thomas J. McCaffrey Carpenter Steel Division Carpenter Technology Corporation J. McClain Danville Division Wyman-Gordon Company T.K. McCluhan Elkem Metals Company D.B. McCutcheon Steltech Technical Services Ltd. Hal L. Miller Nelson Wire Company K.L. Miller The Timken Company Frank Minden Lone Star Steel Michael Mitchell Rockwell International R.W. Monroe Steel Founders' Society of America Timothy E. Moss Inland Bar and Structural Division Inland Steel Company Brian Murkey R.B. & W. Corporation T.E. Murphy Inland Bar and Structural Division Inland Steel Company Janet Nash American Iron and Steel Institute Drew V. Nelson Mechanical Engineering Department Stanford University G.B. Olson Northwestern University George H. Osteen Chaparral Steel J. Otter Saginaw Division General Motors Corporation D.E. Overby Stelco Technical Services Ltd. John F. Papp U.S. Bureau of Mines Y.J. Park Amax Research Company D.F. Paulonis


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United Technologies Leander F. Pease III Powder-Tech Associates, Inc. Thoni V. Philip TVP Inc. Thomas A. Phillips Department of the Interior U.S. Bureau of Mines K.E. Pinnow Crucible Research Center Crucible Materials Corporation Arnold Plant Samuel G. Keywell Company Christopher Plummer The WEFA Group J.A. Pojeta LTV Steel Company R. Randall Rariton River Steel P. Repas U.S.S. Technical Center USX Corporation M.K. Repp The Timken Company Richard Rice Battelle Memorial Institute William L. Roberts Consultant G.J. Roe Bethlehem Steel Corporation Kurt Rohrbach Carpenter Technology Corporation A.R. Rosenfield Battelle Memorial Institute James A. Rossow Wyman-Gordon Company C.P. Royer Exxon Production Research Company Mamdouh M. Salama Conoco Inc. Norman L. Samways Association of Iron and Steel Engineers Gregory D. Sander Ring Screw Works J.A. Schmidt Joseph T. Ryerson and Sons, Inc. Michael Schmidt Carpenter Technology Corporation W. Schuld Seneca Wire & Manufacturing Company R.E. Schwer Cannon-Muskegon Corporation Kay M. Shupe Bliss & Laughlin Steel Company V.K. Sikka Oak Ridge National Laboratory Steve Slavonic


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Teledyne Columbia-Summerill Dale L. Smith Argonne National Laboratory Richard B. Smith Western Steel Division Stanadyne, Inc. Dennis Smyth The Algoma Steel Corporation Ltd. G.R. Speich Department of Metallurgical Engineering Illinois Institute of Technology Thomas Spry Commonwealth Edition W. Stasko Crucible Materials Corporation Crucible Research Center Doru M. Stefanescu The University of Alabama Joseph R. Stephens Lewis Research Center National Aeronautics and Space Administration P.A. Stine General Electric Company N.S. Stoloff Rensselaer Polytechnic Institute John R. Stubbles LTV Steel Company D.K. Subramanyam Ergenics, Inc. A.E. Swansiger ABC Rail Corporation R.W. Swindeman Oak Ridge National Laboratory N. Tepovich Connecticut Steel Millicent H. Thomas LTV Steel Company Geoff Tither Niobium Products Company, Inc. George F. Vander Voort Carpenter Technology Corporation Elgin Van Meter Empire-Detroit Steel Division Cyclops Corporation Krishna M. Vedula Materials Science & Engineering Department Case Western Reserve University G.M. Waid The Timken Company Charles F. Walton Consultant Lee R. Walton Latrobe Steel Company Yung-Shih Wang Exxon Production Research Company S.D. Wasko Allegheny Ludlum Steel Division Allegheny Ludlum Corporation


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J.R. Weeks Brookhaven National Laboratory Charles V. White GMI Engineering and Management Institute Alexander D. Wilson Lukens Steel Company Peter H. Wright Chaparral Steel Company B. Yalamanchili North Star Steel Texas Company Z. Zimerman Bethlehem Steel Corporation

Foreword For nearly 70 years the Metals Handbook has been one of the most widely read and respected sources of information on the subject of metals. Launched in 1923 as a single volume, it has remained a durable reference work, with each succeeding edition demonstrating a continuing upward trend in growth, in subject coverage, and in reader acceptance. As we enter the final decade of the 20th century, the ever-quickening pace of modern life has forced an increasing demand for timely and accurate technical information. Such a demand was the impetus for this, the 10th Edition of Metals Handbook. Since the publication of Volume 1 of the 9th Edition in 1978, there have been significant technological advances in the field of metallurgy. The goal of the present volume is to document these advances as they pertain to the properties and selection of cast irons, steels, and superalloys. A companion volume on properties and selection of nonferrous alloys, special-purpose materials, and pure metals will be published this autumn. Projected volumes in the 10th Edition will present expanded coverage on processing and fabrication of metals; testing, inspection, and failure analysis; microstructural analysis and materials characterization; and corrosion and wear phenomena (the latter a subject area new to the Handbook series). During the 12 years it took to complete the 17 volumes of the 9th Edition, the high standards for technical reliability and comprehensiveness for which Metals Handbook is internationally known were retained. Through the collective efforts of the ASM Handbook Committee, the editorial staff of the Handbook, and nearly 200 contributors from industry, research organizations, government establishments, and educational institutions, Volume 1 of the 10th Edition continues this legacy of excellence. Klaus M. Zwilsky President ASM INTERNATIONAL Edward L. Langer Managing Director ASM INTERNATIONAL

Preface During the past decade, tremendous advances have taken place in the field of materials science. Rapid technological growth and development of composite materials, plastics, and ceramics combined with continued improvements in ferrous and nonferrous metals have made materials selection one of the most challenging endeavors for engineers. Yet the process of selection of materials has also evolved. No longer is a mere recitation of specifications, compositions, and properties adequate when dealing with this complex operation. Instead, information is needed that explains the correlation among the processing, structures, and properties of materials as well as their areas of use. It is the aim of this volume⎯the first in the new 10th Edition series of Metals Handbook⎯to present such data. Like the technology it documents, the Metals Handbook is also evolving. To be truly effective and valid as a reference work, each Edition of the Handbook must have its own identity. To merely repeat information, or to simply make superficial cosmetic changes, would be self-defeating. As such, utmost care and thought were brought to the task of planning the 10th Edition by both the ASM Handbook Committee and the Editorial Staff. To ensure that the 10th Edition continued the tradition of quality associated with the Handbook, it was agreed that it was necessary to: • Determine which subjects (articles) not included in previous Handbooks needed to be added to the 10th Edition • Determine which previously published articles needed only to be revised and/or expanded • Determine which previously published articles needed to be completely rewritten


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• Determine which areas needed to be de-emphasized • Identify and eliminate obsolete data The next step was to determine how the subject of properties selection should be addressed in the 10th Edition. Considering the information explosion that has taken place during the past 30 years, the single-volume approach used for Volume 1 of the 8th Edition (published in 1961) was not considered feasible. For the 9th Edition, three separate volumes on properties and selection were published from 1978 to 1980. This approach, however, was considered somewhat fragmented, particularly in regard to steels: carbon and low-alloy steels were covered in Volume 1, whereas tools steels, austenitic manganese steels, and stainless steels were described in Volume 3. After considering the various options, it was decided that the most logical and user-friendly approach would be to publish two comprehensive volumes on properties and selection. In the present volume, emphasis has been placed on cast irons, carbon and low-alloy steels, and high-performance alloys such as stainless steels and superalloys. A companion volume on properties and selection of nonferrous alloys and special-purpose materials will follow (see Table 1 for an abbreviated table of contents). Table 1 Abbreviated table of contents for Volume 2, 10th Edition, Metals Handbook Specific Metals and Alloys ÃWrought Aluminum and Aluminum Alloys ÃCast Aluminum Alloys ÃAluminum-Lithium Alloys ÃAluminum P/M Alloys ÃWrought Copper and Copper Alloys ÃCast Copper Alloys ÃCopper P/M Products ÃNickel and Nickel Alloys ÃBeryllium-Copper and Beryllium-Nickel Alloys ÃCobalt and Cobalt Alloys ÃMagnesium and Magnesium Alloys ÃTin and Tin Alloys ÃZinc and Zinc Alloys ÃLead and Lead Alloys ÃRefractory Metals and Alloys ÃWrought Titanium and Titanium Alloys ÃCast Titanium Alloys ÃTitanium P/M Alloys ÃZirconium and Hafnium ÃUranium and Uranium Alloys ÃBeryllium ÃPrecious Metals ÃRare Earth Metals ÃGermanium and Germanium Compounds ÃGallium and Gallium Compounds ÃIndium and Bismuth Special-Purpose Materials ÃSoft Magnetic Materials ÃPermanent Magnet Materials ÃMetallic Glasses ÃSuperconducting Materials ÃElectrical Resistance Alloys ÃElectric Contact Materials ÃThermocouple Materials ÃLow Expansion Alloys ÃShape-Memory Alloys


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ÃMaterials For Sliding Bearings ÃMetal-Matrix Composite Materials ÃOrdered Intermetallics ÃCemented Carbides ÃCermets ÃSuperabrasives and Ultrahard Tool Materials ÃStructural Ceramics Pure Metals ÃPreparation and Characterization of Pure Metals ÃProperties of Pure Metals Special Engineering Topics ÃRecycling of Nonferrous Alloys ÃToxicity of Metals

Principal Sections Volume 1 has been organized into seven major sections: • • • • • • •

Cast Irons Carbon and Low-Alloy Steels Hardenability of Carbon and Low-Alloy Steels Fabrication Characteristics of Carbon and Low-Alloy Steels Service Characteristics of Carbon and Low-Alloy Steels Specialty Steels and Heat-Resistant Alloys Special Engineering Topics

Of the 53 articles contained in these sections, 14 are new, 10 were completely rewritten, and the remaining articles have been substantially revised. A review of the content of the major sections is given below; highlighted are differences between the present volume and its 9th Edition predecessor. Table 2 summarizes the content of the principal sections. Table 2 Summary of contents for Volume 1, 10th Edition, Metals Handbook Section title Cast Irons

Number of articles

Pages

Figures(a)

Tables(b)

References

6

104

155

81

108

21

344

298

266

230

Hardenability of Carbon and Low-Alloy Steels

3

122

210

178

28

Fabrication Characteristics of Carbon and Low-Alloy Steels

4

44

56

10

85

Carbon and Low-Allow Steels

Service Characteristics of Carbon and Low-Alloy Steels Specialty Steels and Heat-Resistant Alloys Special Engineering Topics

6

140

219

22

567

11

252

249

163

358

2

27

29

11

50

ÃTotals 53 1033 1216 731 1426 (a) Total number of figure captions; some figures may include more than one illustration. (b) Does not include unnumbered in-text tables or tables that are part of figures

Cast irons are described in six articles. The introductory article on "Classification and Basic Metallurgy of Cast Irons" was completely rewritten for the 10th Edition. The article on "Compacted Graphite Iron" is new to the Handbook. Both of these contributions were authored by D.M. Stefanescu (The University of Alabama), who served as Chairman of Volume 15, Casting, of the 9th Edition. The remaining four articles contain new information on materials (for example, austempered ductile iron) and testing (for example, dynamic tear testing). Carbon and Low-Alloy Steels. Key additions to this section include articles that explain the relationships among processing (both melt and rolling processes), microstructures, and properties of steels. Of particular note is the article by G. Krauss (Colorado School of Mines) on pages 126 to 139 and the various articles on high-strength low-alloy steels. Other highlights include an extensive tabular compilation that cross-references SAE-AISI steels to their international counterparts (see the article "Classification and Designation of Steels" ) and an article on "Bearing Steels" that compares both case-hardened and through-hardened bearing materials.


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Hardenability of Carbon and Low-Alloy Steels. Following articles that introduce H-steels and describe hardenability concepts, including test procedures to determine the hardening response of steels, a comprehensive collection of hardenability curves is presented. Both English and metric hardenability curves are provided for some 86 steels. Fabrication Characteristics. Sheet formability, forgeability, machinability, and weldability are described next. The article on bulk formability, which emphasizes recent studies on HSLA forging steels, is new to the Handbook series. The material on weldability was completely rewritten and occupies nearly four times the space allotted in the 9th Edition. Service Characteristics. The influence of various in-service environments on the properties of steels is one of the most widely studied subjects in metallurgy. Among the topics described in this section are elevated-temperature creep properties, low-temperature fracture toughness, fatigue properties, and impact toughness. A new article also describes the deleterious effect of neutron irradiation on alloy and stainless steels. Of critical importance to this section, however, is the definitive treatise on "Embrittlement of Steels" written by G.F. Vander Voort (Carpenter Technology Corporation). Featuring more than 75 graphs and 372 references, this 48-page article explores the causes and effects of both thermal and environmental degradation on a wide variety of steels. Compared with the 9th Edition on the same subject, this represents a nearly tenfold increase in coverage. Specialty Steels and Heat-Resistant Alloys. Eleven articles on wrought, cast, and powder metallurgy materials for specialty and/or high-performance applications make up this section. Alloy development and selection criteria as related to corrosion-resistant and heat-resistant steels and superalloys are well documented. More than 100 pages are devoted to stainless steels, while three new articles have been written on superalloys⎯including one on newly developed directionally solidified and single-crystal nickel-base alloys used for aerospace engine applications. Special Engineering Topics. The final section examines two subjects that are becoming increasingly important to the engineering community: (1) the availability and supply of strategic materials, such as chromium and cobalt, used in stainless steel and superalloy production, and (2) the current efforts to recycle highly alloyed materials. Both of these subjects are new to the Handbook series. A second article on recycling of nonferrous alloys will be published in Volume 2 of the 10th Edition. Acknowledgments Successful completion of this Handbook required the cooperation and talents of literally hundreds of professional men and women. In terms of the book's technical content, we are indebted to the authors, reviewers, and miscellaneous contributors−some 200 strong−upon whose collective experience and knowledge rests the accuracy and authority of the volume. Thanks are also due to the ASM Handbook Committee and its capable Chairman, Dennis D. Huffman (The Timken Company). The ideas and suggestions provided by members of the committee proved invaluable during the two years of planning required for the 10th edition. Lastly, we would like to acknowledge the efforts of those companies who have worked closely with ASM's editorial and production staff on this and many other Handbook volumes. Our thanks go to Byrd Data Imaging for their tireless efforts in maintaining a demanding typesetting schedule, to Rand McNally & company for the care and quality brought to printing the Handbook, and to Precision Graphics, Don O. Tech, Accurate Art, and HaDel Studio for their attention to detail during preparation of Handbook artwork. Their combined efforts have resulted in a significant and lasting contribution to the metals industry. The Editors

General Information Officers and Trustees of ASM INTERNATIONAL (1990−1991) Klaus M. Zwilsky President and Trustee National Materials Advisory Board National Academy of Sciences Stephen M. Copley Vice President and Trustee Illinois Institute of Technology Richard K. Pitler Immediate Past President and Trustee Allegheny Ludlum Corporation (retired) Edward L. Langer Secretary and Managing Director ASM INTERNATIONAL Robert D. Halverstadt Treasurer AIMe Associates


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Trustees John V. Andrews Teledyne Allvac Edward R. Burrell Inco Alloys International, Inc. H. Joseph Klein Haynes International, Inc. Kenneth F. Packer Packer Engineering, Inc. Hans Portisch VDM Technologies Corporation William E. Quist Boeing Commercial Airplanes John G. Simon General Motors Corporation Charles Yaker Howmet Corporation Daniel S. Zamborsky Consultant

Members of the ASM Handbook Committee (1990−1991) Dennis D. Huffman (Chairman 1986−; Member 1983−) The Timken Company Roger J. Austin (1984−) ABARIS Roy G. Baggerly (1987−) Kenworth Truck Company Robert J. Barnhurst (1988−) Noranda Research Centre Hans Borstell (1988−) Grumman Aircraft Systems Gordon Bourland (1988−) LTV Aerospace and Defense Company John F. Breedis (1989−) Olin Corporation Stephen J. Burden (1989−) GTE Valenite Craig V. Darragh (1989−) The Timken Company Gerald P. Fritzke (1988−) Metallurgical Associates J. Ernesto Indacochea (1987−) University of Illinois at Chicago John B. Lambert (1988−) Fansteel Inc. James C. Leslie (1988−) Advanced Composites Products and Technology Eli Levy (1987−) The De Havilland Aircraft Company of Canada William L. Mankins (1989−) Inco Alloys International, Inc. Arnold R. Marder (1987−)


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Lehigh University John E. Masters (1988−) American Cyanamid Company David V. Neff (1986−) Metaullics Systems David LeRoy Olson (1982−1988; 1989−) Colorado School of Mines Dean E. Orr (1988−) Orr Metallurgical Consulting Service, Inc. Edwin L. Rooy (1989−) Aluminum Company of America Kenneth P. Young (1988−) AMAX Research & Development

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) 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) N.E. Promisel (1955−1961) (Member, 1954−1963) G.J. Shubat (1973−1975) (Member, 1966−1975) W.A. Stadtler (1969−1972) (Member, 1962−1972)


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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 Robert L. Stedfeld, Director of Reference Publications, Joseph R. Davis, Manager of Handbook Development; Kathleen M. Mills, Manager of Book Production; Steven R. Lampman, Technical Editor; Theodore B. Zorc, Technical Editor; Heather F. Lampman, Editorial Supervisor; George M. Crankovic, Editorial Coordinator; Alice W. Ronke, Assistant Editor; Scott D. Henry, Assistant Editor; Janice L. Daquila, Assistant Editor; Janet Jakel, Word Processing Specialist; Karen Lynn O'Keefe, Word Processing Specialist. Editorial assistance was provided by Lois A. Abel, Robert T. Kiepura, Penelope Thomas, and Nikki D. Wheaton. Conversion to Electronic Files ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High-Performance Alloys was converted to electronic files in 1997. The conversion was based on the Fourth Printing (1995). 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, Grace Davidson, Randall Boring, Robert Braddock, and Kathleen Dragolich. 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 © 1990 by ASM International All Rights Reserved. Metals 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 Metals 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 Metals 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) Metals Handbook/Prepared under the direction of the ASM International Handbook Committee _10th ed. Includes bibliographies and indexes. Contents: v. 1. Properties and Selection: Irons, Steels, and High-Performance Alloys. 1. Metals⎯Handbooks, manuals, etc. I. ASM International. Handbook Committee. II. Title: ASM Handbook. TA459.M43 1990 620.1'6 90−115 ISBN 0-87170-377-7 (v.1) SAN 204-7586 ISBN 0-87170-380-7 Printed in the United States of America


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Classification and Basic Metallurgy of C...

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Cast Irons Classification and Basic Metallurgy of Cast Iron Doru M. Stefanescu, The University of Alabama THE TERM CAST IRON, like the term steel, identifies a large family of ferrous alloys. Cast irons are multicomponent ferrous alloys, which solidify with a eutectic. They contain major (iron, carbon, silicon), minor (0.1%) elements. Cast iron has higher carbon and silicon contents than steel. Because of the higher carbon content, the structure of cast iron, as opposed to that of steel, exhibits a rich carbon phase. Depending primarily on composition, cooling rate, and melt treatment, cast iron can solidify according to the thermodynamically metastable Fe-Fe3C system or the stable Fe-Gr system. When the metastable path is followed, the rich carbon phase in the eutectic is the iron carbide; when the stable solidification path is followed, the rich carbon phase is graphite. Referring only to the binary Fe-Fe3C or Fe-Gr system, cast iron can be defined as an iron-carbon alloy with more than 2%C. The reader is cautioned that silicon and other alloying elements may considerably change the maximum solubility of carbon in austenite (γ). Therefore, in exceptional cases, alloys with less than 2% C can solidify with a eutectic structure and therefore still belong to the family of cast iron. The formation of stable or metastable eutectic is a function of many factors including the nucleation potential of the liquid, chemical composition, and cooling rate. The first two factors determine the graphitization potential of the iron. A high graphitization potential will result in irons with graphite as the rich carbon phase, while a low graphitization potential will result in irons with iron carbide. A schematic of the structure of the common types of commercial cast irons, as well as the processing required to obtain them, is shown in Fig. 1 . Fig. 1 Basic microstructures and processing for obtaining common commercial cast irons


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The two basic types of eutectics⎯the stable austenite-graphite or the metastable austenite-iron carbide (Fe3C)⎯have wide differences in their mechanical properties, such as strength, hardness, toughness, and ductility. Therefore, the basic scope of the metallurgical processing of cast iron is to manipulate the type, amount, and morphology of the eutectic in order to achieve the desired mechanical properties.

Classification Historically, the first classification of cast iron was based on its fracture. Two types of iron were initially recognized: • White iron: Exhibits a white, crystalline fracture surface because fracture occurs along the iron carbide plates; it is the result of metastable solidification (Fe3C eutectic) • Gray iron: Exhibits a gray fracture surface because fracture occurs along the graphite plates (flakes); it is the result of stable solidification (Gr eutectic) With the advent of metallography, and as the body of knowledge pertinent to cast iron increased, other classifications based on microstructural features became possible: • Graphite shape: Lamellar (flake) graphite (FG), spheroidal (nodular) graphite (SG), compacted (vermicular) graphite (CG), and temper graphite (TG); temper graphite results from a solid-state reaction (malleabilization) • Matrix: Ferritic, pearlitic, austenitic, martensitic, bainitic (austempered) This classification is seldom used by the floor foundryman. The most widely used terminology is the commercial one. A first division can be made in two categories: • Common cast irons: For general-purpose applications, they are unalloyed or low alloy


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• Special cast irons: For special applications, generally high alloy The correspondence between commercial and microstructural classification, as well as the final processing stage in obtaining common cast irons, is given in Table 1 . A classification of cast irons by their commercial names and structure is also given in the article "Classification of Ferrous Casting Alloys" in Casting, Volume 15 of ASM Handbook, formerly 9th Edition Metals Handbook. Table 1 Classification of cast iron by commercial designation, microstructure, and fracture Commercial designation

Carbon-rich phase

Matrix(a)

Fracture

Final structure after

Gray iron

Lamellar graphite

P

Gray

Solidification

Ductile iron

Spheroidal graphite

F, P, A

Silver-gray

Solidification or heat treatment

Compacted graphite iron

Compacted vermicular graphite

F, P

Gray

Solidification

White iron

Fe3C

P, M

White

Solidification and heat treatment(b)

Mottled iron

Lamellar Gr + Fe3C

P

Mottled

Solidification

Malleable iron

Temper graphite

F, P

Silver-gray

Heat treatment

Austempered ductile iron Spheroidal graphite At Silver-gray Heat treatment (a) F, ferrite; P, pearlite; A, austenite; M, martensite; At, austempered (bainite). (b) White irons are not usually heat treated, except for stress relief and to continue austenite transformation.

Special cast irons differ from the common cast irons mainly in the higher content of alloying elements (>3%), which promote microstructures having special properties for elevated-temperature applications, corrosion resistance, and wear resistance. A classification of the main types of special cast irons is shown in Fig. 2 . Fig. 2 Classification of special high-alloy cast irons. Source: Ref 1

Principles of the Metallurgy of Cast Iron The goal of the metallurgist is to design a process that will produce a structure that will yield the expected mechanical properties. This requires knowledge of the structure-properties correlation for the particular alloy under consideration as well as of the factors affecting the structure. When discussing the metallurgy of cast iron, the main factors of influence on the structure that one needs to address are: • Chemical composition • Cooling rate • Liquid treatment


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• Heat treatment In addition, the following aspects of combined carbon in cast irons should also be considered: • In the original cooling or through subsequent heat treatment, a matrix can be internally decarburized or carburized by depositing graphite on existing sites or by dissolving carbon from them • Depending on the silicon content and the cooling rate, the pearlite in iron can vary in carbon content. This is a ternary system, and the carbon content of pearlite can be as low as 0.50% with 2.5% Si • The conventionally measured hardness of graphitic irons is influenced by the graphite, especially in gray iron. Martensite microhardness may be as high as 66 HRC, but measures as low as 54 HRC conventionally in gray iron (58 HRC in ductile) • The critical temperature of iron is influenced (raised) by silicon content, not carbon content The following sections in this article discuss some of the basic principles of cast iron metallurgy. More detailed descriptions of the metallurgy of cast irons are available in separate articles in this Volume describing certain types of cast irons. The Section "Ferrous Casting Alloys" in Casting, Volume 15 of ASM Handbook, formerly 9th Edition Metals Handbook, also contains more detailed descriptions on the metallurgy of cast irons. Gray Iron (Flake Graphite Iron) The composition of gray iron must be selected in such a way as to satisfy three basic structural requirements: • The required graphite shape and distribution • The carbide-free (chill-free) structure • The required matrix For common cast iron, the main elements of the chemical composition are carbon and silicon. Figure 3 shows the range of carbon and silicon for common cast irons as compared with steel. It is apparent that irons have carbon in excess of the maximum solubility of carbon in austenite, which is shown by the lower dashed line. A high carbon content increases the amount of graphite or Fe3C. High carbon and silicon contents increase the graphitization potential of the iron as well as its castability. Fig. 3 Carbon and silicon composition ranges of common cast irons and steel. Source: Ref 2

The combined influence of carbon and silicon on the structure is usually taken into account by the carbon equivalent (CE): CE = % C + 0.3(% Si) Ã+ 0.33(% P) − 0.027(% Mn) + 0.4(% S) (Eq 1) Additional information on carbon equivalent is available in the article "Thermodynamic Properties of Iron-Base Alloys" in Casting, Volume 15 of ASM Handbook, formerly 9th Edition Metals Handbook. Although increasing the carbon and silicon contents improves the graphitization potential and therefore decreases the chilling tendency, the strength is adversely affected (Fig. 4 ). This is due to ferrite promotion and the coarsening of pearlite. Fig. 4 General influence of carbon equivalent on the tensile strength of gray iron. Source: Ref 2


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The manganese content varies as a function of the desired matrix. Typically, it can be as low as 0.1% for ferritic irons and as high as 1.2% for pearlitic irons, because manganese is a strong pearlite promoter. From the minor elements, phosphorus and sulfur are the most common and are always present in the composition. They can be as high as 0.15% for low-quality iron and are considerably less for high-quality iron, such as ductile iron or compacted graphite iron. The effect of sulfur must be balanced by the effect of manganese. Without manganese in the iron, undesired iron sulfide (FeS) will form at grain boundaries. If the sulfur content is balanced by manganese, manganese sulfide (MnS) will form, which is harmless because it is distributed within the grains. The optimum ratio between manganese and sulfur for an FeS-free structure and maximum amount of ferrite is: % Mn = 1.7(% S) + 0.15 (Eq 2) Other minor elements, such as aluminum, antimony, arsenic, bismuth, lead, magnesium, cerium, and calcium, can significantly alter both the graphite morphology and the microstructure of the matrix. The range of composition for typical unalloyed common cast irons is given in Table 2 . The typical composition range for lowand high-grade unalloyed gray iron (flake graphite iron) cast in sand molds is given in Table 3 . Table 2 Range of compositions for typical unalloyed common cast irons Composition, % Type of iron

C

Si

Mn

P

S

Gray (FG)

2.5−4.0

1.0−3.0

0.2−1.0

0.002−1.0

0.02−0.25

Compacted graphite (CG)

2.5−4.0

1.0−3.0

0.2−1.0

0.01−0.1

0.01−0.03

Ductile (SG)

3.0−4.0

1.8−2.8

0.1−1.0

0.01−0.1

0.01−0.03

White

1.8−3.6

0.5−1.9

0.25−0.8

0.06−0.2

0.06−0.2

2.2−2.9

0.9−1.9

0.15−1.2

0.02−0.2

0.02−0.2

Malleable (TG) Source: Ref 2

Table 3 Compositions of unalloyed gray irons ASTM A 48 class

Composition, %

Carbon equivalent

C

20B

4.5

3.1−3.4

55B

3.6

≤3.1

Si

Mn

P

S

2.5−2.8

0.5−0.7

0.9

0.15

1.4−1.6

0.6−0.75

0.1

0.12

Both major and minor elements have a direct influence on the morphology of flake graphite. The typical graphite shapes for flake graphite are shown in Fig. 5 . Type A graphite is found in inoculated irons cooled with moderate rates. In general, it is associated with the best mechanical properties, and cast irons with this type of graphite exhibit moderate undercooling during solidification (Fig. 6 ). Type B graphite is found in irons of near-eutectic composition, solidifying on a limited number of nuclei. Large eutectic cell size and low undercoolings are common in cast irons exhibiting this type of graphite. Type C graphite occurs in hypereutectic irons as a result of solidification with minimum undercooling. Type D graphite is found in hypoeutectic or eutectic irons solidified at rather high cooling rates, while type E graphite is characteristic for strongly hypoeutectic irons. Types D and E are both associated with high undercoolings during solidification. Not only graphite shape but also graphite size is important, because it is directly related to strength (Fig. 7 ). Fig. 5 Typical flake graphite shapes specified in ASTM A 247. A, uniform distribution, random orientation; B, rosette groupings; C, kish graphite (superimposed flake sizes, random orientation); D, interdendritic segregation with random orientation; E, interdendritic segregation with preferred orientation


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Fig. 6 Characteristic cooling curves associated with different flake graphite shapes. TE, equilibrium eutectic temperature

Fig. 7 Effect of maximum graphite flake length on the tensile strength of gray iron. Source: Ref 3

Alloying elements can be added in common cast iron to enhance some mechanical properties. They influence both the graphitization potential and the structure and properties of the matrix. The main elements are listed below in terms of their graphitization potential: High positive graphitization potential (decreasing positive potential from top to bottom) Carbon Tin Phosphorus Silicon Aluminum Copper Nickel


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Neutral Iron High negative graphitization potential (increasing negative potential from top to bottom) Manganese Chromium Molybdenum Vanadium

This classification is based on the thermodynamic analysis of the influence of a third element on carbon solubility in the Fe-C-X system, where X is a third element (see the section "Influence of a Third Element on Carbon Solubility in the Fe-C-X System" in the article "Thermodynamic Properties of Iron-Base Alloys" in Casting, Volume 15 of ASM Handbook, formerly 9th Edition Metals Handbook. Although listed as a graphitizer (which may be true thermodynamically), phosphorus also acts as a matrix hardener. Above its solubility level (probably about 0.08%), phosphorus forms a very hard ternary eutectic. The above classification should also include sulfur as a carbide former, although manganese and sulfur can combine and neutralize each other. The resultant manganese sulfide also acts as nuclei for flake graphite. In industrial processes, nucleation phenomena may sometimes override solubility considerations. In general, alloying elements can be classified into three categories. Each is discussed below. Silicon and aluminum increase the graphitization potential for both the eutectic and eutectoid transformations and increase the number of graphite particles. They form solid solutions in the matrix. Because they increase the ferrite/pearlite ratio, they lower strength and hardness. Nickel, copper, and tin increase the graphitization potential during the eutectic transformation, but decrease it during the eutectoid transformation, thus raising the pearlite/ferrite ratio. This second effect is due to the retardation of carbon diffusion. These elements form solid solution in the matrix. Because they increase the amount of pearlite, they raise strength and hardness. Chromium, molybdenum, tungsten, and vanadium decrease the graphitization potential at both stages. Thus, they increase the amount of carbides and pearlite. They concentrate in principal in the carbides, forming (FeX)nC-type carbides, but also alloy the αFe solid solution. As long as carbide formation does not occur, these elements increase strength and hardness. Above a certain level, any of these elements will determine the solidification of a structure with both Gr and Fe3C (mottled structure), which will have lower strength but higher hardness. In alloyed gray iron, the typical ranges for the elements discussed above are as follows: Element Chromium

Composition, % 0.2−0.6

Molybdenum Vanadium

0.2−1 0.1−0.2

Nickel

0.6−1

Copper

0.5−1.5

Tin

0.04−0.08

The influence of composition and cooling rate on tensile strength can be estimated using (Ref 3): TS = 162.37 + 16.61/D − 21.78(% C) Ã−61.29(% Si) − 10.59 (% Mn − 1.7% S) Ã+ 13.80(% Cr) + 2.05(% Ni) + 30.66(% Cu) Ã+ 39.75(% Mo) + 14.16 (% Si)2 Ã−26.25(% Cu)2 − 23.83 (% Mo)2 (Eq 3) where D is the bar diameter (in inches). Equation 3 is valid for bar diameters of 20 to 50 mm (1=8to 2 in.) and compositions within the following ranges: Element

Composition, %

Carbon

3.04−3.29

Chromium

0.1−0.55

Molybdenum

0.03−0.78

Silicon

1.6−2.46

Nickel

0.07−1.62

Sulfur

0.089−0.106

Manganese


0.39−0.98

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Copper

01 Sep 2005

0.07−0.85

The cooling rate, like the chemical composition, can significantly influence the as-cast structure and therefore the mechanical properties. The cooling rate of a casting is primarily a function of its section size. The dependence of structure and properties on section size is termed section sensitivity. Increasing the cooling rate will: • Refine both graphite size and matrix structure; this will result in increased strength and hardness • Increase the chilling tendency; this may result in higher hardness, but will decrease the strength Consequently, composition must be tailored in such a way as to provide the correct graphitization potential for a given cooling rate. For a given chemical composition and as the section thickness increases, the graphite becomes coarser, and the pearlite/ferrite ratio decreases, which results in lower strength and hardness (Fig. 8 ). Higher carbon equivalent has similar effects. Fig. 8 Influence of section thickness of the casting on tensile strength (a) and hardness (b) for a series of gray irons classified by their strength as-cast in 30 mm (1.2 in.) diam bars. Source: Ref 2

The liquid treatment of cast iron is of paramount importance in the processing of this alloy because it can dramatically change the nucleation and growth conditions during solidification. As a result, graphite morphology, and therefore properties, can be significantly affected. In gray iron practice, the liquid treatment used is termed inoculation and consists of minute additions of minor elements before pouring. Typically, ferrosilicon with additions of aluminum and calcium, or proprietary alloys are used as inoculants. The main effects of inoculation are: • An increased graphitization potential because of decreased undercooling during solidification; as a result of this, the chilling tendency is diminished, and graphite shape changes from type D or E to type A • A finer structure, that is, higher number of eutectic cells, with a subsequent increase in strength As shown in Fig. 9 , inoculation improves tensile strength. This influence is more pronounced for low-CE cast irons. Fig. 9 Influence of inoculation on tensile strength as a function of carbon equivalent for 30 mm (1.2 in.) diam bars. Source: Ref 2


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01 Sep 2005

Heat treatment can considerably alter the matrix structure, although graphite shape and size remain basically unaffected. A rather low proportion of the total gray iron produced is heat treated. Common heat treatment may consist of stress relieving or of annealing to decrease hardness. Ductile Iron (Spheroidal Graphite Iron) Composition. The main effects of chemical composition are similar to those described for gray iron, with quantitative differences in the extent of these effects and qualitative differences in the influence on graphite morphology. The carbon equivalent has only a mild influence on the properties and structure of ductile iron, because it affects graphite shape considerably less than in the case of gray iron. Nevertheless, to prevent excessive shrinkage, high chilling tendency, graphite flotation, or a high impact transition temperature, optimum amounts of carbon and silicon must be selected.Figure 10 shows the basic guidelines for the selection of appropriate compositions. Fig. 10 Typical range for carbon and silicon contents in good-quality ductile iron. Source: Ref 2

As mentioned previously, minor elements can significantly alter the structure in terms of graphite morphology, chilling tendency, and matrix structure. Minor elements can promote the spheroidization of graphite or can have an adverse effect on graphite shape. The minor elements that adversely affect graphite shape are said to degenerate graphite shape. A variety of graphite shapes can occur, as illustrated in Fig. 11 . Graphite shape is the single most important factor affecting the mechanical properties of cast iron, as shown in Fig. 12 . Fig. 11 Typical graphite shapes after ASTM A 247. I, spheroidal graphite; II, imperfect spheroidal graphite; III, temper graphite, IV, compacted graphite; V, crab graphite; VI, exploded graphite; VII, flake graphite


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Fig. 12 Influence of graphite morphology on the stress-strain curve of several cast irons

The generic influence of various elements on graphite shape is given in Table 4 . The elements in the first group⎯the spheroidizing elements⎯can change graphite shape from flake through compacted to spheroidal. This is illustrated in Fig. 13 for magnesium. The most widely used element for the production of spheroidal graphite is magnesium. The amount of residual magnesium, Mgresid, required to produce spheroidal graphite is generally 0.03 to 0.05%. The precise level depends on the cooling rate. A higher cooling rate requires less magnesium. The amount of magnesium to be added in the iron is a function of the initial


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01 Sep 2005

sulfur level, Sin, and the recovery of magnesium, η, in the particular process used:

(Eq 4) Table 4 Influence of minor elements on graphite shape Element category

Element

Spheroidizer

Magnesium, calcium, rare earths (cerium, lanthanum, etc.), yttrium

Neutral

Iron, carbon, alloying, elements

Antispheroidizer (degenerate shape)

Aluminum, arsenic, bismuth, tellurium, titanium, lead, sulfur, antimony

Fig. 13 Influence of residual magnesium on graphite shape

A residual magnesium level that is too low results in insufficient nodularity (that is, a low ratio between the spheroidal graphite and the total amount of graphite in the structure). This in turn results in a deterioration of the mechanical properties of the iron, as illustrated in Fig. 14 . If the magnesium content is too high, carbides are promoted. Fig. 14 Influence of residual magnesium (a) and nodularity (b) on some mechanical properties of ductile iron. Sources: Ref 4, 5


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01 Sep 2005

The presence of antispheroidizing (deleterious) minor elements may result in graphite shape deterioration, up to complete graphite degeneration. Therefore, upper limits are set on the amount of deleterious elements to be accepted in the composition of cast iron. Typical limits are given below (Ref 6): Element Aluminum

Composition, % 0.1

Arsenic

0.02

Bismuth

0.002

Cadmium

0.01

Lead

0.002

Antimony

0.002

Selenium

0.03

Tellurium

0.02

Titanium

0.1

Zirconium

0.1

These values can be influenced by the combination of various elements and by the presence of rare earths in the composition. Furthermore, some of these elements can be deliberately added during liquid processing in order to increase nodule count. Alloying elements have in principle the same influence on structure and properties as for gray iron. Because a better graphite morphology allows more efficient use of the mechanical properties of the matrix, alloying is more common in ductile iron than in gray iron. Cooling Rate. When changing the cooling rate, effects similar to those discussed for gray iron also occur in ductile iron, but the section sensitivity of ductile iron is lower. This is because spheroidal graphite is less affected by cooling rate than flake graphite. The liquid treatment of ductile iron is more complex than that of gray iron. The two stages for the liquid treatment of ductile iron are: • Modification, which consists of magnesium or magnesium alloy treatment of the melt, with the purpose of changing graphite shape from flake to spheroidal • Inoculation (normally, postinoculation, that is, after the magnesium treatment) to increase the nodule count. Increasing the nodule count is an important goal, because a higher nodule count is associated with less chilling tendency (Fig. 15 ) and a higher as-cast ferrite/pearlite ratio


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01 Sep 2005

Fig. 15 Influence of the amount of 75% ferrosilicon added as a postinoculant on the nodule count and chill depth of 3 mm (0.12 in.) plates. Source: Ref 7

Heat treatment is extensively used in the processing of ductile iron because better advantage can be taken of the matrix structure than for gray iron. The heat treatments usually applied are as follows: • • • • •

Stress relieving Annealing to produce a ferritic matrix Normalizing to produce a pearlitic matrix Hardening to produce tempering structures Austempering to produce a ferritic bainite

The advantage of austempering is that it results in ductile irons with twice the tensile strength for the same toughness. A comparison between some mechanical properties of austempered ductile iron and standard ductile iron is shown in Fig. 16 . Fig. 16 Properties of some standard and austempered ductile irons. Source: Ref 8

Compacted Graphite Irons Compacted graphite irons have a graphite shape intermediate between spheroidal and flake. Typically, compacted graphite looks like type IV graphite (Fig. 11 ). Consequently, most of the properties of CG irons lie in between those of gray and ductile iron. The chemical composition effects are similar to those described for ductile iron. Carbon equivalent influences strength less obviously than for the case of gray iron, but more than for ductile iron, as shown in Fig. 17 . The graphite shape is controlled, as in the case of ductile iron, through the content of minor elements. When the goal is to produce compacted graphite, it is easier from the standpoint of controlling the structure to combine spheroidizing (magnesium, calcium, and/or rare earths) and antispheroidizing (titanium and/or aluminum) elements. Additional information is available in the article"Compacted Graphite Irons" in Casting, Volume 15 of ASM Handbook, formerly 9th Edition Metals Handbook. Fig. 17 Effect of carbon equivalent on the tensile strength of flake, compacted, and spheroidal graphite irons cast in 30


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01 Sep 2005

mm (1.2 in.) diam bars. Source: Ref 9

The cooling rate affects properties less for gray iron but more for ductile iron (Fig. 18 ). In other words, CG iron is less section sensitive than gray iron. However, high cooling rates are to be avoided because of the high propensity of CG iron for chilling and high nodule count in thin sections. Fig. 18 Influence of section thickness on the tensile strength of CG irons. Source: Ref 10

Liquid treatment can have two stages, as for ductile iron. Modification can be achieved with magnesium, Mg + Ti, Ce + Ca, and so on. Inoculation must be kept at a low level to avoid excessive nodularity. Heat treatment is not common for CG irons. Malleable Irons Malleable cast irons differ from the types of irons previously discussed in that they have an initial as-cast white structure, that is, a structure consisting of iron carbides in a pearlitic matrix. This white structure is then heat treated (annealing at 800 to 970 °C, or 1470 to 1780 °F), which results in the decomposition of Fe3C and the formation of temper graphite. The basic solid state reaction is: Fe3C → γ + Gr (Eq 5) The final structure consists of graphite and pearlite, pearlite and ferrite, or ferrite. The structure of the matrix is a function of the cooling rate after annealing. Most of the malleable iron is produced by this technique and is called blackheart malleable iron. Some malleable iron is produced in Europe by decarburization of the white as-cast iron, and it is called whiteheart malleable iron. The composition of malleable irons must be selected in such a way as to produce a white as-cast structure and to allow for fast annealing times. Some typical compositions are given in Table 2 . Although higher carbon and silicon reduce the heat treatment time, they must be limited to ensure a graphite-free structure upon solidification. Both tensile strength and elongation decrease with higher carbon equivalent. Nevertheless, it is not enough to control the carbon equivalent. The annealing time depends on the number of graphite nuclei available for graphitization, which in turn depends on, among other factors, the C/Si ratio. As shown in Fig. 19 , a lower C/Si ratio (that is, a higher silicon content for a constant carbon equivalent) results in a higher temper graphite count. This in turn translates into shorter annealing times. Fig. 19 Influence of C/Si ratio on the number of temper graphite clusters at constant carbon equivalent. Source: Ref 10


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01 Sep 2005

Manganese content and the Mn/S ratio must be closely controlled. In general, a lower manganese content is used when ferritic rather than pearlitic structures are desired. The correct Mn/S ratio can be calculated with Eq 2 . Equation 2 is plotted in Fig. 20 . Under the line described by Eq 2 , all sulfur is stoichiometrically tied to manganese as MnS. The excess manganese is dissolved in the ferrite. In the range delimited by the lines given by Eq 2 and the line Mn/S = 1, a mixed sulfide, (Mn,Fe)S, is formed. For Mn/S ratios smaller than 1, pure FeS is also formed. It is assumed that the degree of compacting of temper graphite depends on the type of sulfides occurring in the iron (Ref 11). When FeS is predominant, very compacted, nodular temper graphite forms, but some undissolved Fe3C may persist in the structure, resulting in lower elongations. When MnS is predominant, although the graphite is less compacted, elongation is higher because of the completely Fe3C-free structure. Fig. 20 Influence of the Mn/S ratio on the shape of temper graphite. Bracketed elements are dissolved in the matrix.

The Mn/S ratio also influences the number of temper graphite particles. From this standpoint, the optimum Mn/S ratio is about 2 to 4 (Fig. 21 ). Fig. 21 Influence of the Mn/S ratio on the number of temper graphite clusters after annealing. A, low-temperature holding for 12 h at 350 °C (660 °F); B, no low-temperature holding

Alloying elements can be used in some grades of pearlitic malleable irons. The manganese content can be increased to 1.2%, or copper, nickel, and/or molybdenum can be added. Chromium must be avoided because it produces stable carbides, which are difficult to decompose during annealing. Cooling Rate. Like all other irons, malleable irons are sensitive to cooling rate. Nevertheless, because the final structure is the result of a solid-state reaction, they are the least section sensitive irons. Typical correlations between tensile strength, elongation, and section thickness are shown in Fig. 22 .


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01 Sep 2005

Fig. 22 Influence of bar diameter on the tensile strength (a) and elongation (b) of blackheart malleable iron. Source: Ref 13

The liquid treatment of malleable iron increases the number of nuclei available for the solid-state graphitization reaction. This can be achieved in two different ways, as follows: • By adding elements that increase undercooling during solidification. Typical elements in this category are magnesium, cerium, bismuth, and tellurium. Higher undercooling results in finer structure, which in turn means more γ-Fe3C interface. Because graphite nucleates at the γ-Fe3C interface, this means more nucleation sites for graphite. Higher undercooling during solidification also prevents the formation of unwanted eutectic graphite • By adding nitrite-forming elements to the melt. Typical elements in this category are aluminum, boron, titanium, and zirconium The heat treatment of malleable iron determines the final structure of this iron. It has two basic stages. In the first stage, the iron carbide is decomposed in austenite and graphite (Eq 5 ). In the second stage, the austenite is transformed into pearlite, ferrite, or a mixture of the two. Although there are some compositional differences between ferritic and pearlitic irons, the main difference is in the heat treatment cycle. When ferritic structures are to be produced, cooling rates in the range of 3 to 10 °C/h (5 to 18 °F/h) are required through the eutectoid transformation in the second stage. This is necessary to allow for a complete austenite-to-ferrite reaction. A typical annealing cycle for ferritic malleable iron is shown in Fig. 23 . When pearlitic irons are to be produced, different schemes can be used, as shown in Fig. 24 . The goal of the treatment is to achieve a eutectoid transformation according to the austenite-to-pearlite reaction. In some limited cases, quenching-tempering treatments are used for malleable irons. Fig. 23 Heat treatment cycle for ferritic blackheart malleable iron. Source: Ref 1

Fig. 24 Heat treatment cycles for pearlitic blackheart malleable irons


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01 Sep 2005

Special Cast Irons Special cast irons, as previously discussed, are alloy irons that take advantage of the radical changes in structure produced by rather large amounts of alloying elements. Abrasion resistance can be improved by increasing hardness, which in turn can be achieved by either increasing the amount of carbides and their hardness or by producing a martensitic structure. The least expensive material is white iron with a pearlitic matrix. Additions of 3 to 5% Ni and 1.5 to 2.5% Cr result in irons with (FeCr)3C carbides and an as-cast martensitic matrix. Additions of 11 to 35% Cr produce (CrFe)7C3 carbides, which are harder than the iron carbides. Additions of 4 to 16% Mn will result in a structure consisting of (FeMn) 3C, martensite, and work-hardenable austenite. Heat resistance depends on the stability of the microstructure. Irons used for these applications may have a ferritic structure with graphite (5% Si), a ferritic structure with stable carbides (11 to 28% Cr), or a stable austenitic structure with either spheroidal or flake graphite (18% Ni, 5% Si). For corrosion resistance, irons with high chromium (up to 28%), nickel (up to 18%), and silicon (up to 15%) are used. REFERENCES 1. R. Elliot, Cast Iron Technology, Butterworths, 1988 2. C.F. Walton and T.J. Opar, Ed., Iron Castings Handbook, Iron Castings Society, 1981 3. C.E. Bates, AFS Trans., Vol 94, 1986, p 889 4. R. Barton, B.C.I.R.A.J., No. 5, 1961, p 668 5. R.W. Lindsay and A. Shames, AFS Trans., Vol 60, 1952, p 650 6. H. Morrogh, AFS Trans., Vol 60, 1952, p 439 7. D.M. Stefanescu, AFS Int. Cast Met. J., June 1981, p 23 8. J.F. Janowak and R.B. Gundlach, AFS Trans., Vol 91, 1983, p 377 9. G.F. Sergeant and E.R. Evans, Br. Foundryman, May 1978, p 115 10. D.M. Stefanescu, Metalurgia, No. 7, 1967, p 368 11. K. Roesch, Stahl Eisen, No. 24, 1957, p 1747 12. R.P. Todorov, in Proceedings of the 32nd International Foundry Congress (Warsaw, Poland), International Committee of Foundry Technical Associations 13. K.M. Ankab, O.E. Shulte, and P.N. Bidulia, Isvestia Vishih Utchebnik Zavedenia-Tchornaia, Metallurghia, No. 5, 1966, p 168


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Gray Iron

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Cast Irons Gray Iron Revised by Charles V. White, GMI Engineering and Management Institute CAST IRONS are alloys of iron, carbon, and silicon in which more carbon is present than can be retained in solid solution in austenite at the eutectic temperature. In gray cast iron, the carbon that exceeds the solubility in austenite precipitates as flake graphite. Gray irons usually contain 2.5 to 4% C, 1 to 3% Si, and additions of manganese, depending on the desired microstructure (as low as 0.1% Mn in ferritic gray irons and as high as 1.2% in pearlitics). Sulfur and phosphorus are also present in small amounts as residual impurities. Certain important but low-tonnage specialty items in this family of cast metals (notably the austenitic and other highly alloyed gray irons) are not dealt with here; instead the emphasis is on the properties of gray irons used most often and in the largest tonnages. Information on the high-alloy gray irons is given in the article "Alloy Cast Irons" in this Volume. The basic metallurgy of gray cast irons is discussed in the article "Classification and Basic Metallurgy of Cast Iron" in this Volume.

Classes of Gray Iron A simple and convenient classification of the gray irons is found in ASTM specification A 48, which classifies the various types in terms of tensile strength, expressed in ksi. The ASTM classification by no means connotes a scale of ascending superiority from class 20 (minimum tensile strength of 140 MPa, or 20 ksi) to class 60 (minimum tensile strength of 410 MPa, or 60 ksi). In many applications strength is not the major criterion for the choice of grade. For example, for parts such as clutch plates and brake drums, where resistance to heat checking is important, low-strength grades of iron are the superior performers. Similarly, in heat shock applications such as ingot or pig molds, a class 60 iron would fail quickly, whereas good performance is shown by class 25 iron. In machine tools and other parts subject to vibration, the better damping capacity of low-strength irons is often advantageous. Generally, it can be assumed that the following properties of gray cast irons increase with increasing tensile strength from class 20 to class 60: • • • •

All strengths, including strength at elevated temperature Ability to be machined to a fine finish Modulus of elasticity Wear resistance

On the other hand, the following properties decrease with increasing tensile strength, so that low-strength irons often perform better than high-strength irons when these properties are important: • • • •

Machinability Resistance to thermal shock Damping capacity Ability to be cast in thin sections

Applications Gray iron is used for many different types of parts in a very wide variety of machines and structures. Like parts made from other metals and alloys, parts intended to be produced as gray iron castings must be evaluated for the specific service conditions before being approved for production. Often a stress analysis of prototype castings helps establish the appropriate class of gray iron as well as any proof test requirements or other acceptance criteria for production parts.

Castability


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Gray Iron

01 Sep 2005

Successful production of a gray iron casting depends on the fluidity of the molten metal and on the cooling rate (which is influenced by the minimum section thickness and on section thickness variations). Casting design is often described in terms of section sensitivity. This is an attempt to correlate properties in critical sections of the casting with the combined effects of composition and cooling rate. All these factors are interrelated and may be condensed into a single term, castability, which for gray iron may be defined as the minimum section thickness that can be produced in a mold cavity with given volume/area ratio and mechanical properties consistent with the type of iron being poured. Fluidity. Scrap losses resulting from misruns, cold shuts, and round corners are often attributed to the lack of fluidity of the metal being poured. Mold conditions, pouring rate, and other process variables being equal, the fluidity of commercial gray irons depends primarily on the amount of superheat above the freezing temperature (liquidus). As the total carbon (TC) content decreases, the liquidus temperature increases, and the fluidity at a given pouring temperature therefore decreases. Fluidity is commonly measured as the length of flow into a spiral-type fluidity test mold. The relation between fluidity and superheat is shown in Fig. 1 for four unalloyed gray irons of different carbon contents. Fig. 1 Fluidity versus degree of superheat for four gray irons of different carbon contents

The significance of the relationships between fluidity, carbon content, and pouring temperature becomes apparent when it is realized that the gradation in strength in the ASTM classification of gray iron is due in large part to differences in carbon content (~3.60 to 3.80% for class 20; ~2.70 to 2.95% for class 60). The fluidity of these irons thus resolves into a measure of the practical limits of maximum pouring temperature as opposed to the liquidus of the iron being poured. These practical limits of maximum pouring temperature are largely determined by three factors: • The ability of both mold and cores to withstand the impact of molten iron, an ability that decreases as the pouring temperature increases, thereby favoring low pouring temperatures • The fact that metal tap temperatures seldom exceed 1550 °C (2825 °F). Because ladling and reladling to the point of pouring generally accounts for temperature losses of 55 to 85 °C (100 to 150 °F), the final pouring temperatures seldom exceed 1450 to 1495 °C (2640 to 2720 °F), and in most instances maximum pouring temperatures in the range 1410 to 1450 °C (2570 to 2640 °F) are considered more realistic • The necessity to control the overall thermal input to the mold in order to control the final desired microstructure It can be seen from Table 1 that because of differences in liquidus temperature, the amount of superheat (and therefore fluidity) varies with carbon content when various compositions are cast from the same pouring temperature. Table 1 Superheat above liquidus for 2% Si irons of various carbon contents poured at 1455 °C (2650 °F) Liquidus temperature

Superheat above liquidus

°C

°F

°C

°F

2.52

1295

2360

160

290

3.04

1245

2270

210

380

3.60

1175

2150

280

500

Carbon, %


Page 35


Gray Iron

01 Sep 2005

Microstructure The usual microstructure of gray iron is a matrix of pearlite with graphite flakes dispersed throughout. Foundry practice can be varied so that nucleation and growth of graphite flakes occur in a pattern that enhances the desired properties. The amount, size, and distribution of graphite are important. Cooling that is too rapid may produce so-called chilled iron, in which the excess carbon is found in the form of massive carbides. Cooling at intermediate rates can produce mottled iron, in which carbon is present in the form of both primary cementite (iron carbide) and graphite. Very slow cooling of irons that contain large percentages of silicon and carbon is likely to produce considerable ferrite and pearlite throughout the matrix, together with coarse graphite flakes. Flake graphite is one of seven types (shapes or forms) of graphite established in ASTM A 247. Flake graphite is subdivided into five types (patterns), which are designated by the letters A through E (see Fig. 2 ). Graphite size is established by comparison with an ASTM size chart, which shows the typical appearances of flakes of eight different sizes at 100× magnification. Fig. 2 Types of graphite flakes in gray iron (AFS-ASTM). In the recommended practice (ASTM A 247), these charts are shown at a magnification of 100×. They have been reduced to one-third size for reproduction here.

Type A flake graphite (random orientation) is preferred for most applications. In the intermediate flake sizes, type A flake graphite is superior to other types in certain wear applications such as the cylinders of internal combustion engines. Type B flake graphite (rosette pattern) is typical of fairly rapid cooling, such as is common with moderately thin sections (about 10 mm, or 3 =8in.) and along the surfaces of thicker sections, and sometimes results from poor inoculation. The large flakes of type C flake graphite are typical of kish graphite that is formed in hypereutectic irons. These large flakes enhance resistance to thermal shock by increasing thermal conductivity and decreasing elastic modulus. On the other hand, large flakes are not conducive to good surface finishes on machined parts or to high strength or good impact resistance. The small, randomly oriented interdendritic flakes in type D flake graphite promote a fine machined finish by minimizing surface pitting, but it is difficult to obtain a pearlitic matrix with this type of graphite. Type D flake graphite may be formed near rapidly cooled surfaces or in thin sections . Frequently, such graphite is surrounded by a ferrite matrix, resulting in soft spots in the casting. Type E flake graphite is an interdendritic form, which has a preferred rather than a random orientation. Unlike type D graphite, type E graphite can be associated with a pearlitic matrix and thus can produce a casting whose wear properties are as good as those of a casting containing only type A graphite in a pearlitic matrix. There are, of course, many applications in which flake type has no significance as long as the mechanical property requirements are met. Solidification of Gray Iron. In a hypereutectic gray iron, solidification begins with the precipitation of kish graphite in the melt. Kish grows as large, straight, undistorted flakes or as very thick, lumpy flakes that tend to rise to the surface of the melt because of their low relative density. When the temperature has been lowered sufficiently, the remaining liquid solidifies as a eutectic structure of austenite and graphite. Generally, eutectic graphite is finer than kish graphite. In hypoeutectic iron, solidification begins with the formation of proeutectic austenite dendrites. As the temperature falls, the dendrites grow, and the carbon content of the remaining liquid increases. When the increasing carbon content and decreasing temperature reach eutectic values, eutectic solidification begins. Eutectic growth from many different nuclei proceeds along crystallization fronts that are approximately spherical. Ultimately, the eutectic cells meet and consume the liquid remaining in the spaces between them. During eutectic solidification, the austenite in the eutectic becomes continuous with the dendritic proeutectic austenite, and the structure can be described as a dispersion of graphite flakes in austenite. After solidification, the eutectic cell structure and the proeutectic austenite dendrites cannot be distinguished metallographically except by special etching or in strongly hypoeutectic iron. With eutectic compositions, obviously, solidification takes place as the molten alloy is cooled through the normal eutectic temperature range, but without the prior formation of a proeutectic constituent. During the solidification process, the controlling factor remains the rate at which the solidification is proceeding. The rapid solidification favored by thin section sizes or highly conductive molding media can result in undercooling. Undercooling can cause the solidification to start at a temperature lower than the expected eutectic temperature for a given composition (Fig. 3 ). This can result in a modification of the carbon form from A to E type or can completely suppress its formation and form primary carbides instead. Fig. 3 Undercooling from rapid cooling of a eutectic composition


Page 36


Gray Iron

01 Sep 2005

Room-Temperature Structure. Upon cooling from the eutectic temperature, the austenite will decompose, first by precipitating some of the dissolved carbon and then, at the eutectoid temperature, by undergoing complete transformation. The actual products of the eutectoid transformation depend on rate of cooling as well as on composition of the austenite, but under normal conditions the austenite will transform either to pearlite or to ferrite plus graphite. Transformation to ferrite plus graphite is most likely to occur with slow cooling rates, which allow more time for carbon migration within the austenite; high silicon contents, which favor the formation of graphite rather than cementite; high values of carbon equivalent; and the presence of fine undercooled (type D) flake graphite. Graphite formed during decomposition is deposited on the existing graphite flakes. When carbon equivalent values are relatively low or when cooling rates are relatively fast, the transformation to pearlite is favored. In some instances, the microstructure will contain all three constituents: ferrite, pearlite, and graphite. With certain compositions, especially alloy gray irons, it is possible to produce a martensitic matrix by oil quenching through the eutectoid transformation range or an austempered matrix by appropriate isothermal treatment (Ref 1). These treatments are often done deliberately in a secondary heat treatment where high strength or hardness is especially desired, such as in certain wear applications. The secondary heat treatment of gray iron castings is of great value in producing components that must be hard when machining requirements prohibit the use of components that are cast to final shape in white iron.

Section Sensitivity In practice, the minimum thickness of section in which any given class of gray iron may be poured is more likely to depend on the cooling rate of the section than on the fluidity of the metal. For example, although a plate 300 mm (12 in.) square by 6 mm (0.24 in.) thick can be poured in class 50 as well as in class 25 iron, the former casting would not be gray iron because the cooling rate would be so rapid that massive carbides would be formed. Yet it is entirely feasible to use class 50 iron for a diesel engine cylinder head that has predominantly 6 mm (0.24 in.) wall sections in the water jackets above the firing deck. This is simply because the cooling rate of the cylinder head is reduced by the "mass effect" resulting from enclosed cores and the proximity (often less than 12 mm, or 0.47 in.) of one 6 mm (0.24 in.) wall to the other. Thus the shape of the casting has an important bearing on the choice of metal specification. It should be recognized that the smallest section that can be cast gray, without massive carbides, depends not only on metal composition, but also on foundry practices. For example, by adjusting silicon content or by using graphitizing additions called inoculants in the ladle, the foundryman can decrease the minimum section size for freedom from carbides for a given basic composition of gray iron. The mass effect associated with increasing section thickness or decreasing cooling rate is much more pronounced in gray iron than in cast steel. The mass effect in cast steel results in increased grain size in heavy sections. This also applies to gray iron, but the most important effects are on graphite size and distribution, and on amount of combined carbon. For any given gray iron composition, the rate of cooling from the freezing temperature to below about 650 °C (1200 °F) determines the ratio of combined to graphitic carbon, which controls the hardness and strength of the iron. For this reason the effect of section size in gray iron is considerably greater than in the more homogeneous ferrous metals in which cooling rate does not affect the form and distribution of carbon throughout the metal structure. Typical Effects of Section Size. When a wedge-shape bar with about a 10° taper is cast in a sand mold and sectioned near the center of the length, and Rockwell hardness determinations are made on the cut surface from the point of the wedge progressively into the thicker sections, the curves so determined show to what extent continually increasing section size affects hardness (Fig. 4 ). Fig. 4 Effect of section thickness on hardness and structure. Hardness readings were taken at increasing distance from the tip of a cast wedge section, as shown by inset. Composition of iron: 3.52% C, 2.55% Si, 1.01% Mn, 0.215% P, and 0.086% S. Source: Ref 2


Page 37


Gray Iron

01 Sep 2005

Progressing along the curve from the left in Fig. 4 , the following metallographic constituents occur. The tip of the wedge is white iron (a mixture of carbide and pearlite) with a hardness greater than 50 HRC. As the iron becomes mottled (a mixture of white iron and gray iron), the hardness decreases sharply. A minimum is reached because of the occurrence of fine type D flake graphite, which usually has associated ferrite in large amounts. With a slightly lower cooling rate, the structure becomes fine type A flake graphite in a pearlite matrix with the hardness rising to another maximum on the curve. This structure is usually the most desirable for wear resistance and strength. With increasing section thickness beyond this point, the graphite flakes become coarser, and the pearlite lamellae become more widely spaced, resulting in slightly lower hardness. With further increase in wedge thickness and decrease in cooling rate, pearlite decomposes progressively to a mixture of ferrite and graphite, resulting in softer and weaker iron. The structures of most commercial gray iron castings are represented by the right-hand downward-sloping portion of the curve in Fig. 4 , beyond 5 mm (0.2 in.) wedge thickness, and increasing section size is normally reflected by the gradual lowering of hardness and strength. However, thin sections may be represented by the left-hand downward-sloping portion. Figure 5 shows the average tensile strength (up to ten tests per point) of two irons, for each of which six sizes of cylindrical round bars were cast and appropriate tensile specimens machined. With the class 20 iron, strength increases as the as-cast section decreases down to the 6 mm (0.24 in.) cast bar. However, for the class 30 iron, a section 6 mm (0.24 in.) in diameter is so small that the strength falls off sharply, because of the occurrence of type D flake graphite or mottled iron, or both. The other graph in Fig. 5 shows similar data for the same two classes of iron and for three higher classes. Fig. 5 Effect of section diameter on tensile strength at center of cast specimen for five classes of gray iron

Section sensitivity effects are used in the form of a wedge test in production control to judge the suitability of an iron for pouring a particular casting. In this test, a wedge-shape casting is poured and upon solidification is evaluated. The standard W2 wedge block specified in ASTM A 367 is shown in Fig. 6 . The evaluation consists of measuring the length of the "chilled zone." The measurement, usually made in 0.8 mm (1=32in.) increments, is related to empirically determined data obtained from a "good" casting. If the evaluation indicates an excessive sensitivity for a part, corrections are made to the molten metal prior to pouring. Fig. 6 Standard W2 wedge block used for measuring depth of chill (ASTM A 367). Dimensions given in inches


Page 38


Gray Iron

01 Sep 2005

Volume/Area Ratios. It is extremely difficult to predict with accuracy the cooling rate for castings other than fairly simple shapes. However, because minimum limitations are involved here, the problem can be resolved through comparisons of the casting design with ratios of volume to surface area or with minimum plate sections. The volume/area (V/A) ratios for round, square, and plate sections provide a fairly accurate indication of the minimum casting sections possible in simple geometrical shapes (Table 2 ). The V/A ratios can be reported in either English or metric units and can be converted simply by treating them as length measurements. Table 2 Volume/area (V/A) ratios for round bars, square bars, and plates V/A ratio Cast form and size

mm

in.

Bar, 13 mm ( =2 in.) diam × 533 mm (21 in.)

3.1

0.12

1

Bar, 13 mm ( =2 in.) square × 533 mm (21 in.)

3.1

0.12

Plate, 6.4 × 305 × 305 mm ( =4 × 12 × 12 in.)

3.0

0.12

Bar, 30 mm (1.2 in.) diam × 533 mm (21 in.) (a)

7.4

0.29

Bar, 30 mm (1.2 in.) square × 533 mm (21 in.)

7.4

0.29

Plate, 16 × 305 × 305 mm ( =8 × 12 × 12 in.)

7.1

0.28

Bar, 50 mm (2 in.) diam × 560 mm (22 in.)

12.2

0.48

Bar, 50 mm (2 in.) square × 560 mm (22 in.)

12.2

0.48

Plate, 28.5 × 305 × 305 mm (1 =8 × 12 × 12 in.)

12.0

0.47


22.9

0.90


22.9

0.90

Plate, 65 × 305 × 305 mm (2 =16 × 12 × 12 in.)

22.8

0.90


32.7

1.29


32.7

1.29

Plate, 114 × 305 × 305 mm (4 =2 × 12 × 12 in.) (a) ASTM size B test bar Source: Ref 3

32.7

1.29

1

1

5

1

9

1

Comparison of the ratios of volume to surface area for different shapes gives good agreement with the actual cooling rates of castings made in the same mold material. For long round bars and infinite flat plates, V/A is diameter/4 for bars and thickness/2 for plates; that is, a large plate casting would have the same cooling rate as a round bar with a diameter twice the plate thickness. Most castings, however, freeze somewhat faster than an infinite flat plate, and rather than establishing a 2-to-1 ratio of bar to plate, a smaller ratio will often give a better correlation with the cooling rate. The bar and plate sizes shown in Table 3 are nearly equivalent in cooling rate. Table 3 Bar and plate sizes of equivalent cooling rate For 305 mm (12 in.) square plates, as recorded in Table 2 Bar diameter, in. 1 1 =2 =4

Plate thickness, in.

Ratio of bar diameter to plate thickness 2.0

1.2

5

=8

1.92

2

11=8

1.78

4

29=16

1.56


Page 39


6

Gray Iron

01 Sep 2005

41=2

1.33

Similar comparisons have been made for production castings. In one study, the properties of a flat section from a 0.6 m (24 in.) cross pipe fitting having a nominal thickness of 29.5 mm (1.16 in.) were compared with the properties of a 50 mm (2 in.) diam cylindrical test bar cast from the same heat. The tensile strengths of the test bars were within about 16 MPa (2.3 ksi) of the tensile strengths of the cross pipe fittings for eight heats ranging in strength from about 205 to 310 MPa (30 to 45 ksi), an average variation of less than 8%. These results from production castings correlate well with the calculated equivalence given in Table 3 . Other examples of this type of correlation are given in Ref3. Relationships developed for various specific castings are valid when an iron of controlled composition, and therefore of similar section sensitivity, is used consistently. For instance, with a copper-molybdenum iron of well-controlled composition, a tensile strength of 450 MPa (65 ksi) in the ASTM B test specimen has been found to ensure 345 MPa (50 ksi) tensile strength in a cast crankshaft 2.13 m (7 ft) long with sections thicker than 30.5 mm (1.2 in.). Such translation of properties of a small test bar to properties expected in a larger section cannot be done indiscriminately, because different irons may vary widely in section sensitivity.

Prevailing Sections Although the ASTM size B test bar (30.5 mm, or 1.2 in., diam) is the bar most commonly used for all gray irons from class 20 to class 60, ASTM specification A 48 provides a series of bar sizes from which one that approximates the cooling rate in the critical section of the casting can be selected. In practice, it is customary to be somewhat more definite regarding the prevailing values of minimum casting section considered feasible for the various ASTM classes of cast iron. As summarized in Table 4 , these minimum, prevailing sections include the requirement of freedom from carbidic areas. In a platelike section, occasional thinner walls (such as ribs) are of no importance unless they are very thin or are appended to the outer edges of the casting. Table 4 Minimum prevailing casting sections recommended for gray irons ASTM A48 class 20

Minimum thickness in. 1 =8

V/A ratio(a)

mm

in.

mm

3.2

0.06

1.5

25

1

6.4

0.12

3.0

30

3

9.5

0.17

4.3

35

3

9.5

0.17

4.3

40

5

15.9

0.28

7.1

50

3

=4

19.0

0.33

8.4

1

25.4

0.42

10.7

60 (a) V/A ratios are for square plates.

=4 =8 =8 =8

Mechanical properties of class 30 and class 50 gray irons in various sections are shown in Fig. 7 . For class 30 iron, the combined carbon content and hardness are still at a safe level in sections equivalent to a 10 mm (0.4 in.) plate, which has a V/A ratio of about 5 mm (0.20 in.). For class 50 iron, however, both combined carbon and Brinell hardness show marked increases when the thickness of the equivalent plate section is decreased to about 15 mm (0.6 in.), with V/A ratio around 7 mm (0.27 in.). These results are consistent with the minimum prevailing casting sections recommended in Table 4 . Fig. 7 Mechanical properties of class 30 and class 50 gray iron as a function of section size. Composition of the class 30 iron: 3.40% C, 2.38% Si, 0.71% Mn, 0.423% P, and 0.152% S; for the class 50 iron: 2.96% C, 1.63% Si, 1.05% Mn, 0.67% Mo, 0.114% P, and 0.072% S. Source: Ref 4


Page 40


Gray Iron

01 Sep 2005

The hazards involved in pouring a given class of gray iron in a plate section thinner than recommended are discovered when the casting is machined. Typical losses as a result of specifying too high a strength for a prevailing section of 9.5 mm (3=8in.) are given below (rejections were for "hard spots" that made it impossible to machine the castings by normal methods): Class

Rejections, %

35

Negligible

45

25

55

80−100

In marginal applications, a higher class of iron may sometimes be used if the casting is cooled slowly (in effect, increasing the section thickness) by judicious placement of flow-offs and risers. An example is the successful production of a 25 mm (1 in.) diam single-throw crankshaft for an air compressor. This shaft was hard at the extreme ends when poured in class 50 iron. The difficulty was corrected by flowing metal through each end into flow-off risers that adequately balanced the cooling rate at the ends with the cooling rate at the center. In sum, the selection of a suitable grade of gray iron for a specific casting necessarily requires an evaluation of the size and shape of the casting as related to its cooling rate, or volume/area ratio. For a majority of parts, this evaluation need be no more than a determination of whether or not the V/A ratio of the casting exceeds the minimum V/A ratio indicated for the grade considered.

Test Bar Properties Mechanical property values obtained from test bars are sometimes the only available guides to the mechanical properties of the metal in production castings. When test bars and castings are poured from metal of the same chemical history, correlations can be drawn between the thermal history of the casting and that of the test bar. The strength of the test bar gives a relative strength of the casting, corrected for the cooling rate of the various section thickness. Through careful analysis of the critical sections of a casting, accurate predictions of mechanical behavior can be achieved. Usual Tests. Tension and transverse tests on bars that are cast specifically for such tests are the most common methods used for evaluating the strength of gray iron. Yield strength, elongation, and reduction of area are seldom determined for gray iron in standard tension tests. The transverse test measures strength in bending and has the additional advantage that a deflection value may be obtained readily. Minimum specification values are given in Table 5 . Data can usually be obtained faster from the transverse test than from the tension test because machining of the specimen is unnecessary. The surface condition of the bar will affect the transverse test but not the tension test made on a machined specimen. Conversely, the presence of coarse graphite in the center of the bar, which can occur in an iron that is very section sensitive, will affect the tension test but not the transverse test. Table 5 Transverse breaking loads of gray irons tested per ASTM A 438 Corrected transverse breaking load(a)


Page 41


Gray Iron

Approximate tensile strength

01 Sep 2005

A bar(b)

B bar(c)

C bar(d)

ASTM class(a)

MPa

ksi

kg

lb

kg

lb

kg

lb

20

138

20

408

900

816

1800

2720

6,000

25

172

25

465

1025

907

2000

3080

6,800

30

207

30

522

1150

998

2200

3450

7,600

35

241

35

578

1275

1089

2400

3760

8,300

40

276

40

635

1400

1179

2600

4130

9,100

45

310

45

699

1540

1270

2800

4400

9,700

50

345

50

760

1675

1361

3000

4670

10,300

60 414 60 873 1925 1542 3400 5670 12,500 (a) For separately cast test specimens produced in accordance with ASTM A 48, ASTM A 278, ASME SA278, FED QQ-1-652, or any other specification that designates ASME A438 as the test method. Included in specifications only by agreement between producer and purchaser. (b) 22.4 mm (0.88 in.) diam; 305 mm (12 in.) between supports. (c) 30.5 mm (1.20 in.) diam; 460 mm (18 in.) between supports. (d) 50.8 mm (2.00 in.) diam; 610 mm (24 in.) between supports.

Hardness tests, on either test bars or castings, are used as an approximate measure of strength and sometimes as an indication of relative machinability. Relationships between Brinell hardness and tensile strength generally follow the pattern reproduced in Fig. 8 , which shows the variation of tensile strength with Brinell hardness for a series of gray irons produced by a single foundry. The data in Fig. 8 are from ASTM size A and B test bars poured in a series of inoculated gray irons. The successful use of Brinell hardness as a measure of strength depends on whether it can be proved suitable for the application, which may involve service tests or mechanical property tests on specimens cut from production castings. Fig. 8 Relationship between tensile strength and Brinell hardness for a series of inoculated gray irons from a single foundry. Open circles represent unalloyed gray iron, and closed circles represent alloy gray iron. Source: Ref 5

Testing Precautions. In the assessment of mechanical properties for a series of heats, precautions should be taken to ensure minimum statistical variation between bars. By its nature gray iron behaves as a brittle material in tension, with no measurable elongation after fracture. This characteristic can be exaggerated by imposing a nonaxial load during tensile testing, resulting in statistical variations, which may not be a true measure of the quality of the iron. To overcome this tendency, many shops use self-aligning nonthreaded grips in the performance of tensile tests on gray iron tests bars. Typical Specifications. ASTM A 48 is typical of specifications based on test bars. In practice, one of three different standard sizes of separately cast test bars is used to evaluate the properties in the controlling section of the castings. After manufacturer and purchaser agree on a controlling section of the casting, the size of test bar that corresponds, approximately, to the cooling rate expected in that section is designated by a letter (see Table 6 ). Table 6 Test bars designed to match controlling sections of castings (ASTM A 48) Controlling section

Diameter of as-cast test bar

in.

mm

Test bar

mm

in.

50 S (a) All dimensions of test bar by agreement between manufacturer and purchaser

Most gray iron castings for general engineering use are specified as either class 25, 30, or 35. Specification A 48 is based entirely on mechanical properties, and the composition that provides the required properties can be selected by the individual producer. A manufacturer whose major production is medium-section castings of class 35 iron will find, for heavy-section castings for which the 50 mm (2 in.) test bar is required for qualifying, that the same composition will not meet the requirements for class 35. It will qualify only for some lower class, such as 25 or 30. As the thickness of the controlling section increases, the composition must be adjusted to maintain the same tensile strength. SAE standard J431c for gray cast irons (see Tables 7 , 8 , 9 , and 10 ) describes requirements that are more specific than those described in ASTM A 48. An iron intended for heavy sections, such as grade G3500, is specified to have higher strength and hardness in the standard test bar than does grade G2500, which is intended for light-section castings. Typical applications for the various grades are summarized in Table 11 . Table 7 Mechanical properties of SAE J431 automotive gray cast irons Properties determined from as-cast test bar B (30.5 mm, or 1.2 in., diam) Minimum transverse load SAE grade

Minimum deflection

Minimum tensile strength

Hardness, HB

kg

lb

mm

in.

MPa

ksi

G1800

187 max

780

1720

3.6

0.14

124

18

G2500

170−229

910

2000

4.3

0.17

173

25

G3000

187−241

1000

2200

5.1

0.20

207

30

G3500

207−255

1110

2450

6.1

0.24

241

35

G4000

217−269

1180

2600

6.9

0.27

276

40

Table 8 Typical base compositions of SAE J431 automotive gray cast irons See Table 7 for mechanical properties. If either carbon or silicon is on the high side of the range, the other should be on the low side. Composition, % UNS

SAE grade

TC(a)

Mn

Si

P

S

F10004

G1800(b)

3.40−3.70

0.50−0.80

2.80−2.30

0.15

0.15

F10005

G2500(b)

3.20−3.50

0.60−0.90

2.40−2.00

0.12

0.15

F10006

G3000(c)

3.10−3.40

0.60−0.90

2.30−1.90

0.10

0.15

F10007

G3500(c)

0.60−0.90

2.20−1.80

0.08

0.15

F10008 G4000(c) 3.00−3.30 0.70−1.00 (a) TC, total carbon. (b) Ferritic-pearlitic microstructure. (c) Pearlitic microstructure

2.10−1.80

0.07

0.15

3.00−3.30

Table 9 Mechanical properties of SAE J431 automotive gray cast irons for heavy-duty service Properties determined from as-cast test bar B (30.5 mm, or 1.2 in., diam) Minimum transverse load SAE grade

Minimum deflection

Minimum tensile strength

Hardness, HB

kg

lb

mm

in.

MPa

ksi

G2500a

170−229

910

2000

4.3

0.17

173

25

G3500b

207−255

1090

2400

6.1

0.24

241

35

G3500c

207−255

1090

2400

6.1

0.24

241

35

1180

2600

6.9

0.27

276

40

G4000d 241−321(a) (a) Determined on a specified bearing surface.

Table 10 Typical base compositions of SAE J431 automotive gray cast irons for heavy-duty service See Table 9 for mechanical properties. Composition, %(a) UNS

SAE grade

TC

Mn

Si

P

S

F10009

G2500a(b)

3.40 min

0.60−0.90

1.60−2.10

0.12

0.12

F10010

G3500b(c)

3.40 min

1.30−1.80

0.80

0.12


0.60−0.90

Page 43


F10011

Gray Iron

G3500c(c)

3.50 min

01 Sep 2005

0.60−0.90

0.80

1.30−1.80

0.12

F10012 G4000d (d) 0.07 0.12 3.10−3.60 0.60−0.90 1.95−2.40 (a) If either carbon or silicon is on the high side of the range, the other should be on the low side. Alloying elements not listed in this table may be required. (b) Microstructure: size 2 to 4 type A graphite in a matrix of lamellar pearlite containing not more than 15% free ferrite. (c) Microstructure: size 3 to 5 type A graphite in a matrix of lamellar pearlite containing not more than 5% free ferrite or free carbide. (d) Alloy gray iron containing 0.85 to 1.25% Cr, 0.40 to 0.60% Mo, and 0.20 to 0.45% Ni or as agreed. Microstructure: primary carbides and size 4 to 7 type A or E graphite in a matrix of fine pearlite, as determined in a zone at least 3.2 mm (1=8in.) deep at a specified location on a cam surface.

Table 11 Automotive applications of gray cast iron Grade

Typical uses

G1800

Miscellaneous soft iron castings (as-cast or annealed) in which strength is not a primary consideration

G2500

Small cylinder blocks, cylinder heads, air-cooled cylinders, pistons, clutch plates, oil pump bodies, transmission cases, gearboxes, clutch housings, and light-duty brake drums

G2500a

Brake drums and clutch plates for moderate service requirements, where high-carbon iron is desired to minimize heat checking

G3000

Automobile and diesel cylinder blocks, cylinder heads, flywheels, differential carrier castings, pistons, medium-duty brake drums and clutch plates

G3500

Diesel engine blocks, truck and tractor cylinder blocks and heads, heavy flywheels, tractor transmission cases, heavy gearboxes

G3500b

Brake drums and clutch plates for heavy duty service where both resistance to heat checking and higher strength are definite requirements

G3500c

Brake drums for extra heavy duty service

G4000

Diesel engine castings, liners, cylinders, and pistons

G4000d

Camshafts

ASTM specifications other than A 48 include A 159 (automotive), A 126 (valves, flanges, and pipe fittings), A 74 (soil pipe and fittings), A 278 (pressure-containing parts for temperatures up to 340 °C, or 650 °F), A 319 (nonpressure-containing parts for elevated-temperature service) and A 436 (austenitic gray irons for heat, corrosion, and wear resistance). The austenitic gray irons are described in the article "Alloy Cast Irons" in this Volume. ASTM A 438 describes the standard method for performing transverse bending tests on separately cast, cylindrical test bars of gray cast iron. Compressive Strength. When gray iron is used for structural applications such as machinery foundations or supports, the engineer, who is usually designing, to support weight only, bases his calculations on the compressive strength of the material. Table 12 , which summarizes typical values for mechanical properties of the various grades, shows the high compressive strength of gray irons. Figure 9 compares the stress-strain curves in tension and compression for a class 20 and a class 40 gray iron. The compressive strength of gray iron is typically three to four times that of the tensile strength. If loads other than dead weights are involved (unless these loads are constant), the problem is one of dynamic stresses, which requires the consideration of fatigue and damping characteristics. Table 12 Typical mechanical properties of as-cast standard gray iron test bars Tensile strength

Torsional shear strength

Compressive strength

Reversed bending fatigue limit

Transverse load on test bar B

ASTM A 48 class

MPa

ksi

MPa

ksi

MPa

ksi

MPa

ksi

kg

lb

Hardness , HB

20

152

22

179

26

572

83

69

10

839

1850

156

25

179

26

220

32

669

97

79

11.5

987

2175

174

30

214

31

276

40

752

109

97

14

1145

2525

210

35

252

36.5

334

48.5

855

124

110

16

1293

2850

212

40

293

42.5

393

57

965

140

128

18.5

1440

3175

235

50

362

52.5

503

73

1130

164

148

21.5

1630

3600

262

60

431

62.5

610

88.5

1293

187.5

169

24.5

1678

3700

302

Fig. 9 A comparison of stress-strain curves in tension and compression for a class 20 and a class 40 gray iron. Source: Ref 6


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Gray Iron

01 Sep 2005

Tensile strength is considered in selecting gray iron for parts intended for static loads in direct tension or bending. Such parts include pressure vessels, autoclaves housings and other enclosures, valves, fittings, and levers. Depending on the uncertainty of loading, safety factors of 2 to 12 have been used in figuring allowable design stresses. Transverse Strength and Deflection. When an ASTM arbitration bar is loaded as a simple beam and the load and deflection required to break it are determined, the resulting value is converted into a nominal index of strength by using the standard beam formula. The value that is determined is arbitrarily called the modulus of rupture. The values for modulus of rupture are useful for production control, but cannot be used in the design of castings without further analysis and interpretation. Rarely does a casting have a shape such that those areas subject to bending stress have a direct relationship to the round arbitration bar. A more rational approach is to use the tensile strength (or fatigue limit) and, after determining the section modulus of the actual shape, apply the proper bending formula. However, because the difficulty of obtaining meaningful values of tensile strength in tests of small specimens, the load computed in this manner is usually somewhat lower than the actual load required to rupture the part, unless unfavorable residual stresses are present in the finished part. Elongation of gray iron at fracture is very small (of the order of 0.6%) and hence is seldom reported. The designer cannot use the numerical value of permanent elongation in any quantitative manner. Torsional Shear Strength. As shown in Table 12 , most gray irons have high torsional shear strength. Many grades have torsional strength greater than that of some grades of steel. This characteristic, along with low notch sensitivity, makes gray iron a suitable material for shafting of various types, particularly in the grades of higher tensile strength. Most shafts are subjected to dynamic torsional stresses, and the designer should carefully consider the exact nature of the loads to be encountered. For the


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Gray Iron

01 Sep 2005

higher-strength irons, stress concentration factors associated with changes of shape in the part are important for torque loads as well as for bending and tension loads. Modulus of Elasticity. Typical stress-strain curves for gray iron are shown in Fig. 10 . Gray iron does not obey Hooke's law, and the modulus in tension is usually determined arbitrarily as the slope of the line connecting the origin of the stress-strain curve with the point corresponding to one-fourth the tensile strength (secant modulus). Some engineers use the slope of the stress-strain curve near the origin (tangent modulus). The secant modulus is a conservative value suitable for most engineering work; design loads are seldom as high as one-fourth the tensile strength, and the deviation of the stress-strain curve from linearity is usually less than 0.01% at these loads. However, in the design of certain types of machinery, such as precision equipment, where design stresses are very low, the use of the tangent modulus may represent the actual situation more accurately. Fig. 10 Typical stress-strain curves for three classes of gray iron in tension

The modulus of gray iron (see Table 13 ) varies considerably more than do the moduli for most other metals. Thus, in using observed strain to calculate stress, it is essential to measure the modulus of the particular gray iron specimen being considered. A significant range in modulus values is experienced because of both section size and chemical analysis variations (Fig. 11 ). In addition, the modulus experiences a linear reduction with increasing temperature. The rate of reduction can be reduced through alloy additions (Fig. 12 ). The numerical value of the modulus in torsion is always less than it is in tension, as is the case with steel. Table 13 Typical moduli of elasticity of as-cast standard gray iron test bars ASTM A 48 class

Tensile modulus

Torsional modulus

GPa

106 psi

GPa

106 psi

20

66−97

9.6−14.0

27−39

3.9−5.6

25

79−102

11.5−14.8

32−41

4.6−6.0

30

90−113

13.0−16.4

36−45

5.2−6.6

35

100−119

14.5−17.2

40−48

5.8−6.9

40

110−138

16.0−20.0

44−54

6.4−7.8

50

130−157

18.8−22.8

50−55

7.2−8.0

60

141−162

20.4−23.5

54−59

7.8−8.5

Fig. 11 Interrelationship of mechanical properties, section diameter, carbon equivalent, and liquidus temperature of gray iron. Data are from one foundry and are based on dry sand molding and ferrosilicon inoculation.

Fig. 12 Elastic modulus as a function of temperature in various alloyed gray cast irons. Source: Ref 7


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Gray Iron

01 Sep 2005

Hardness of gray iron, as measured by Brinell or Rockwell testers, is an intermediate value between the hardness of the soft graphite in the iron and that of the harder metallic matrix. Variations in graphite size and distribution will cause wide variations in hardness (particularly Rockwell hardness) even though the hardness of the metallic matrix is constant. To illustrate this effect, the matrix microhardness of five types of hardened iron are compared with Rockwell C measurements on the same iron in Table 14 . Table 14 Influence of graphite type and distribution on the hardness of hardened gray irons Total carbon, %

Conventional hardness, HRC(a)

Matrix hardness, HRC(b)

A

3.06

45.2(c)

61.5

A

3.53

43.1

61.0

A

4.00

32.0

62.0

D

3.30

54.0

62.5

Type of graphite

D 3.60 48.7 60.5 (a) Measured by conventional Rockwell C test. (b) Hardness of matrix, measured with superficial hardness tester and converted to Rockwell C. (c) Although this value was obtained in the specific test cited, it is not typical of gray iron of 3.06% C. Ordinarily the hardness of such iron is 48 to 50 HRC.

If any hardness correlation is to be attempted, the type and amount of graphite in the irons being compared must be constant. Rockwell hardness tests are considered appropriate only for hardened castings (such as camshafts), and even on hardened castings, Brinell tests are preferred. Brinell tests must be used when attempting any strength correlations for unhardened castings.

Fatigue Limit in Reversed Bending Because fatigue limits are expensive to determine, the designer usually has incomplete information on this property. Fatigue life curves at room temperature for a gray iron under completely reversed cycles of bending stress are shown in Fig. 13 (a), in which each point represents the data from one specimen. The effects of temperature on fatigue limit and tensile strength are shown in Fig. 13 (b) and 13 (c), respectively. Fig. 13 Effect of temperature on fatigue behavior and tensile strength of a gray iron (2.84% C, 1.52% Si, 1.05% Mn, 0.07% P, 0.12% S, 0.31% Cr, 0.20% Ni, 0.37% Cu). (a) Reversed bending fatigue life at room temperature. (b) Reversed bending fatigue limit at elevated temperatures. (c)Tensile strength at elevated temperatures. Source: Ref 8


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Gray Iron

01 Sep 2005

Axial loading or torsional loading cycles are frequently encountered in designing parts of cast iron, and in many instances these are not completely reversed loads. Types of regularly repeated stress variation can usually be expressed as a function of a mean stress and a stress range. Wherever possible, the designer should use actual data from the limited information available. Without precisely applicable test data, an estimate of the reversed bending fatigue limit of machined parts may be made by using about 35% of the minimum specified tensile strength of the particular grade of gray iron being considered. This value is probably more conservative than an average of the few data available on the fatigue limit for gray iron. An approximation of the effect of range of stress on fatigue limit may be obtained from diagrams such as Fig. 14 . Tensile strength is plotted on the horizontal axis to represent fracture strength under static load (which corresponds to a 0 stress range). Reversed bending fatigue limit is plotted on the ordinate for 0 mean stress, and the two points are joined by a straight line. The resulting diagram yields a fatigue limit (maximum value of alternating stress) for any value of mean stress. Fig. 14 Diagram showing safe and unsafe fatigue zones for cast iron subjected to ranges of alternating stress superimposed on a mean stress. Example point P shows conditions of tensile (positive) mean stress; P′ shows compressive (negative) mean stress. The safety factor is represented by the ratio of OF to OP or OF′ to OP′. For conditions of constant mean tensile stress, DK/DP is the safety factor.

Few data available are applicable to design problems involving dynamic loading where the stress cycle is predominantly compressive rather than tensile. Some work done on aluminum and steel indicates that for compressive (negative) mean stress, the behavior of these materials could be represented by a horizontal line beginning at the fatigue limit in reversed bending, as indicated in Fig. 14 . Gray iron is probably at least as strong as this for loading cycles resulting in negative mean stress, because it is much stronger in static compression than in static tension. It is therefore a natural assumption that the parallel behavior shown in Fig. 14 is conservative. If, prior to design, the real stress cycle can be predicted with confidence and enough data are available for a reliable S-N diagram for the gray iron proposed, the casting might be dimensioned to obtain a minimum safety factor of two based on fatigue strength. (Some uses may require more conservative or more liberal loading.) The approximate safety factor is best illustrated by point P in Fig. 14 . The safety factor is determined by the distance from the origin to the fatigue limit line along a ray through the cyclic-stress point, divided by the distance from the origin to that point. In Fig. 14 this is OF/OP. On this diagram, point P′ represents a stress cycle having a negative mean stress. In other words, the maximum compressive stress is greater than the tensile stress reached during the loading cycle. In this instance, the safety factor is the distance OF′/OP′. However, this analysis assumes that overloads will increase the mean stress and alternating stress in the same proportion. This may not always be true, particularly in systems with mechanical vibration in which the mean stress may remain constant. For this condition, the vertical line through P would be used; that is, DK/DP would be the factor of safety. Most engineers use diagrams such as Fig. 14 mainly to determine whether a given condition of mean stress and cyclic stress results in a design safe for infinite life. The designer can also determine whether variations in the mean stress and the alternating stress that he anticipates will place his design in the unsafe zone. Usually the data required to analyze a particular set of conditions are obtained experimentally. It is emphasized that the number of cycles of alternating stress implied in Fig. 14 is the number normally used to determine fatigue limits, that is, approximately 10 million. Fewer cycles, as encountered in infrequent overloads, will be safer than indicated by a particular point plotted on a diagram for infinite life. Too few data are available to draw a diagram for less than infinite life. Fatigue Notch Sensitivity. In general, very little allowance need be made for a reduction in fatigue strength caused by notches or abrupt changes of section in gray iron members. The low-strength irons exhibit only a slight reduction in strength in the presence of fillets and holes. That is, the notch sensitivity index approaches 0; in other words, the effective stress concentration factor for these notches approaches 1. This characteristic can be explained by considering the graphite flakes in gray iron to be internal notches. Thus, gray iron can be thought of as a material that is already full of notches and that therefore has little or no sensitivity to the presence of additional notches resulting from design features. The strength-reducing effect of the internal notches is included in the fatigue limit values determined by conventional laboratory tests with smooth bars.


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Gray Iron

01 Sep 2005

High-strength irons usually exhibit greater notch sensitivity, but probably not the full theoretical value represented by the stress concentration factor. Normal stress concentration factors (see Ref 9) are probably suitable for high-strength gray irons.

Pressure Tightness Gray iron castings are used widely in pressure applications such as cylinder blocks, manifolds, pipe and pipe fittings, compressors, and pumps. An important design factor for pressure tightness is uniformity of section. Parts of relatively uniform wall section cast in gray iron are pressure tight against gases as well as liquids. Most trouble with leaking castings is encountered when there are unavoidable, abrupt changes in section. Shrinkage, internal porosity, stress cracking, and other defects are most likely to occur at junctions between heavy and light sections. Watertight castings are considerably less challenging to the foundryman than gastight castings. A slight sponginess or internal porosity at heavy sections will not usually leak water, and even if there is slight seepage, internal rusting will soon plug the passages permanently. For gastightness, however, castings must be quite sound. Lack of pressure tightness in gray iron castings can usually be traced to internal porosity, which also is called internal shrinkage. In gray iron this seems to be a phenomenon distinctly different from the normal solidification shrinkage that often appears on the casting surface as a sink or draw, which can be cured by risering. Internal porosity or shrinkage (very difficult to prevent, even by the use of very heavy risers) is usually associated with poor feeding and lack of directional solidification. It can also be associated with heavy inoculation with calcium to promote a high graphite cell count. On the other hand, the use of strontium-bearing inoculants does not reduce cell size and can help control internal shrinkage in marginally gated castings. Visible internal porosity may appear at centers of mass when the phosphorus content exceeds 0.25%. In critically gated castings, visible internal porosity may appear when the phosphorus content is as low as 0.09%. Chromium and molybdenum accentuate this effect of phosphorus, while nickel has a slight mitigating influence. The effect of phosphorus may be caused in part by the fact that lowering phosphorus content also lowers the effective carbon equivalent. Lowering carbon equivalent by reducing carbon or silicon, or both, instead of phosphorus, might similarly reduce the leakage of pressure castings, but other foundry problems (such as increasing amounts of normal shrinkage) would be encountered. ASTM A 278 for pressure castings requires a carbon equivalent of 3.8% (max), a phosphorus content of 0.25% (max), and a sulfur content of 0.12% (max) for castings to be used above 230 °C (450 °F). In addition to composition control, good overall foundry practice is required for consistently producing pressure-tight castings. Sand properties and gating must be controlled to avoid sand inclusions. Pouring temperature must be adequate for good fluidity, and heavy sections should be fed wherever possible. Mold properties have been found to interact with composition to influence internal soundness. Generally, molds rammed tightly produce the soundest castings because of freedom from mold wall movement during solidification.

Impact Resistance Where high impact resistance is needed, gray iron is not recommended. Gray iron has considerably lower impact strength than cast carbon steel, ductile iron, or malleable iron. However, many gray iron castings need some impact strength to resist breakage in shipment or use. There is incomplete agreement on a standard method of impact testing for cast iron. Two methods that have been used successfully are given in ASTM A 327. Most impact testing of cast iron has been used as a research tool. Most producers of cast iron pressure pipe use a routine pipe impact test as a control. Impact resistance in pipe helps avoid breakage in shipping and handling.

Machinability The machinability of most gray cast iron is superior to that of most other cast irons of equivalent hardness, as well as to that of virtually all steel. The flake graphite introduces discontinuities in the metal matrix, which act as chip breakers. The graphite itself serves as a lubricant for the cutting tool. However, economical cutting depends on more than inherent machinability alone. Often, trouble in machining gray iron can be traced to one or more of several factors: the presence of chill of corners and in light sections, the presence of adhering sand on the surface of the casting, swells, usually the result of soft molds, shifted castings, shrinks, and phases included in the matrix as a result of melting practice. Chill at corners and in light sections is more likely to be encountered with small castings, with higher-strength irons, and with designs that have light sections in the cope (or top) of the mold. Most foundries control iron with a chill test that gives an indication of the tendency of the iron to form white or mottled iron in light sections. The foundryman may treat his iron with a small amount (0.5 to 2.5 kg/Mg, or 1 to 5 lb/ton) of a graphitizing alloy such as calcium-bearing ferrosilicon or other proprietary inoculant, which effectively decreases chill. Inoculation to achieve control of the tendency to chill usually does not result in significant changes in the composition or the physical properties of the iron, although it does produce changes in the mechanical properties. Light sections (5 mm, or 3=16in.) usually cannot be cast in gray iron of higher than class 25. Class 30 iron can be cast in 6 mm 1 ( =4in.) sections. These values are different for different designs, depending on how the casting is made and gated. The important thing to understand is that the cooling rate in the mold at the time of freezing determines whether the iron will be gray, white, or mottled. If the thin section is in the drag or near the gate, the flow of hot metal heats the mold, thereby decreasing the rate of


Page 49


Gray Iron

01 Sep 2005

cooling and enhancing the formation of the gray iron. If chill is encountered, it is generally best to correct the trouble at its origin. It is usually uneconomical to anneal castings to remove chill because, in addition to heat treating for 2 h at 900 °C (1650 °F) for unalloyed irons, recleaning may be required for removing scale. Distortion beyond tolerance often occurs, and there are sacrifices in hardness and strength. Adhering sand usually can be removed by effective cleaning, but sand present as the result of penetration of the iron into the mold wall is extremely difficult to blast clean. This is a foundry defect that is best corrected at the source. Slowing the speed of machining and increasing the rate of feed is the best approach to salvaging castings of this type. Carbide tools are better than high-speed tools for resisting the extreme abrasion. Swells are most troublesome in operations such as broaching and in other setups tooled for high production. The additional metal often places an excessive load on the tool, which may chip or dull but not actually fail until some time after the troublesome parts have been machined. In highly automated machining centers, this casting defect can result in significant statistical variation in dimensions and high scrap rates. Shifted castings are similar to swells in their action on cutting tools. Shifts or swells also may cause excessive tool loading if the locating points are affected. It is important to consider the positions of such locating points when designing the castings and also to avoid indiscriminate grinding of locating points in the foundry cleaning room. Shrinks are not present but can be troublesome when encountered in operations such as drilling. For example, the drill may tend to drift from its intended path to follow the shrink, which offers less resistance to the drill. Sometimes a drill may break because it encounters a region of higher hardness, which often is associated with an area of shrink. Cast iron is the easiest metal to cast without internal shrink. Eutectic freezing is accompanied by expansion due to the precipitation of low-density graphite, which aids in obtaining internal soundness. Machinability rating is complex and is discussed in the article "Machinability of Steels" in this Volume. Criteria such as power per unit volume in unit time are of greatest importance in selecting a machine tool and the size of its motor. Machinability ratings based on tool life under standard test conditions are helpful but are not readily interpreted into the economics of machining. One way to indicate the effect on machinability of changing from one grade of iron to another is shown in Table 15 , which is based on an experiment in which metal was removed from four types of gray iron using single-point carbide tools. For each type of iron, cutting speed was varied until the removal of 3300 cm3 (200 in.3) of iron resulted in a 0.75 mm (0.030 in.) wear land on the tool. The values of cutting speed thus obtained serve as qualitative evaluations of machinability, but not as qualitative indices. Optimum cutting speeds, tool materials, cutting fluids, feeds, and finish requirements must be studied for any given machine tool setup. Tool life is an important factor, because the machine must be stopped to change tools and the tools must be resharpened. Progress has been made in decreasing both machine downtime for tool changes and resharpening cost by the use of solid carbide inserts or bits and by other means. Table 15 Machinability of gray iron Tensile strength Microstructure

Cutting speed(a)

ASTM class

MPa

ksi

Hardness, HB

m/min

ft/min

Acicular iron

50

407

59

263

46

150

Fine pearlite, alloy

40

310

45

225

95

310

Ferrite (annealed)

...

108

15.7

100

293

960

Coarse pearlite, no alloy

35

241

35

195

99

325

(a) Cutting speed at which removal of 3280 cm3 (200 in.3 produced 0.75 mm (0.030 in.) wear land on single-point carbide tools. Source: Ref 10

Annealing. The greatly improved machinability obtainable in gray iron by annealing has been advantageous to automotive and other industries for many years. Annealing is usually of the subcritical type, such as 1 h at 730 to 760 °C (1350 to 1400 °F). Some employ a cycle of heating to 785 to 815 °C (1450 to 1500 °F) and cooling at 22 °C/h (40 °F/h) to about 595 °C (1100 °F). These treatments graphitize the carbide in the pearlite and result in a ferritic matrix. Finding it uneconomical to graphitize primary carbide, most users try to avoid obtaining it. Annealing for improved machinability is most economical when the casting is small and the amount of machining is large. The annealing treatments described result in sacrifices in hardness and strength. A typical class 35 iron will be downgraded to about class 20 in strength by this treatment. In applications in which wear resistance is important, such as cylinder blocks, gray iron is not annealed because of the unsatisfactory performance obtained with a ferritic matrix. Table 16 and Fig. 15 show the changes in hardness, strength, and structure of class 35 gray iron obtained by annealing. Table 16 Effect of annealing on hardness and strength of class 35 gray iron Tensile strength Condition

MPa

ksi

Hardness, HB

As-cast

268

38.9

217

Annealed

165

23.9

131


Page 50


Gray Iron

01 Sep 2005

Composition of iron: 3.30% total C, 2.22% Si, 0.027% P, 0.18% S, 0.61% Mn, 0.03% Cr, 0.03% Ni, 0.14% Cu, Mo nil. Annealing treatment consisted of 1 h at 775 °C (1425 °F), followed by cooling in the furnace to 540 °C (1000 °F).

Fig. 15 Structure of class 35 iron, as-cast (left) and after annealing

Wear Gray iron is used widely for machine components that must resist wear. Different types of iron, however, exhibit great differences in wear characteristics. These differences do not correlate with the commonly measured properties of the iron. In the discussion that follows, the general conclusions are related to metal-to-metal wear during sliding contact under conditions of normal lubrication. Although most of the supporting illustrations are for engine cylinders, the results have wider applicability. The published data on gray iron wear are somewhat inconsistent; accelerated-wear tests often do not correlate with field service experience, nor does field experience in one application necessarily agree with that in another application. In many applications, properties of the surface are at least as important as properties of the metal; for instance, wear resistance may frequently be enhanced by lapping the wear areas. Relative hardness between mating parts also may be important to optimum wear resistance. Frequently, a hardness difference no greater than 10 points on the Brinell scale is considered optimum. For components in sliding contact, such as engine cylinders, valve guides, and latheways, the recognized types of wear are cutting wear, abrasive wear, adhesive wear, and corrosive wear. In well-designed machinery, wear proceeds slowly and consists of combinations of very mild forms of these four types of wear. The predominant characteristic of this so-called normal wear is the development of a glazed surface during break-in. The primary objective in any wear application is to establish conditions that produce and maintain this glaze. A secondary objective is to minimize normal wear. Cutting wear is caused by the mechanical removal of surface metal as a result of surface roughness and is similar to the action of a file. It usually occurs during the breaking in of new parts. Abrasive wear is caused by the cutting action or loose, abrasive particles that get between the contacting faces and act like a lapping compound. Under some circumstances, abrasive particles embedded in one or both of the contacting faces can produce a similar action. Adhesive wear is caused by metal-to-metal contact, resultant welding, and the breaking of those welds. When this happens on a large scale (a process known as galling), the metal is smeared, and severe surface damage results. Even in properly operating equipment, some wear occurs by adhesion on a microscopic scale. An intermediate form of adhesive wear called scuffing is less destructive than galling, but still results in an abnormally high rate of metal removal. In all probability, accelerated-wear tests between clean surfaces (with or without the presence of a lubricant) are predominantly tests of galling or scuffing. Corrosive wear is a special type of wear that combines abrasion or adhesion with the chemical action of the environment. In engine cylinders, it is caused by condensed acidic products of combustion during low-temperature operation. Usually this kind of wear cannot be corrected by modifications in ordinary types of gray iron.

Resistance to Scuffing Assuming reasonable design and operating conditions that minimize cutting and abrasive wear, scuffing or mild galling is the abnormal wear condition most likely to defeat the objective of obtaining and maintaining a glazed surface due to normal wear. Several alloy combinations with widely varied microstructures were tested by Shuck (Ref 11). A brake shoe type of specimen was held against a rotating gray iron drum for 1 h, after which both specimen and drum were checked for weight loss. The conditions of this test indicate that wear was caused primarily by adhesion (scuffing), although cutting wear may have had a contributing effect. The tests covered almost all conceivable microstructures and a wide variety of compositions having hardnesses below 300 HB. The conclusions, which basically agreed with other less comprehensive investigations, were: • Microstructure determines wearing characteristics • As the graphite structure becomes coarser and tends toward type A, scuffing decreases • Interdendritic type D graphite and its associated ferrite give very poor results


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Gray Iron

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• Secondary ferrite associated with random type A graphite is less damaging than that associated with type D • Pearlitic, acicular, or tempered martensitic structures in the same hardness range are equal in wear resistance • For a given type of graphite, as the matrix becomes more pearlitic and harder, wear resistance increases These results were substantiated in scuff tests of cylinder sleeves in a diesel engine. The engine, equipped with the test sleeves, was operated at constant speed, and the horsepower was increased by increments above normal until scuffing occurred. The scuffing horsepower was expressed in a ratio with normal horsepower. Effect of Graphite Structure. To eliminate differences in matrix, four types of gray iron were hardened and tempered at 205 °C (400 °F) to produce a uniform martensitic matrix. The data in Table 17 show that the greater the amount and flake size of graphite, the greater the resistance to scuffing. The changes in casting method in tests 3 and 4 caused no significant change in scuffing. The iron in test 5 is one used formerly in high-performance brake drums to give maximum resistance to scoring. Table 17 Effect of graphite structure on resistance to scuffing Test No.

Type of graphite(a)

1

None (5150 steel)

2

100% type D, centrifugally cast

3

Type A, size 4 to 6, some type B, centrifugally cast

4

Same as 3 except cast in sand mold

Total carbon, %

Resistance to scuffing(b)

...

1.45(d) (a) Different chemical compositions were tested in two of the four types of iron. See Table 19 for compositions. Matrix of all specimens was tempered martensite. (b) Expressed as ratio: horsepower to produce scuffing divided by normal horsepower. (c) All the steel sleeves scuffed below normal horsepower. (d) Maximum available engine horsepower produced no scuffing.

Effect of Matrix Structure. Similar tests, in which the graphite was type D and the matrix was varied, gave the results shown in Table 18 . The combination of type D graphite and ferrite in test 1 is especially poor even though the rest of the structure is pearlite. Pearlitic type A iron with small amounts of free ferrite gave ratios greater than 1.33 in the same tests, showing that the results obtained in test 1 are peculiar to the combination of type D graphite and ferrite. When type D graphite occurs in as-cast irons, it is usually associated with ferrite. Consequently, foundry practices that avoid the formation of type D graphite should be adopted when producing castings for wear applications. Table 18 Effect of matrix microstructure on resistance to scuffing Test No.

Microstructure of matrix

Hardness

Resistance to scuffing(a)

1

Pearlite, with ferrite occurring in graphitic areas

196−227 HB

0.10%

−0.0155 to −0.042

Mg + Ti 0.05 − 0.1% Ã+ Al 0.2 − 0.3%

−0.0110 to −0.055

Mg + Al > 0.35%

−0.0060 to −0.35

(a) ∆S: final % S − 0.34 (% residual elements) − 1.33 (% Mg)


Page 103


Compacted Graphite Iron

01 Sep 2005

Compacted graphite iron has a strong ferritization tendency. Copper, tin, molybdenum, and even aluminum can be used to increase the pearlite/ferrite ratio. Again, the optimum amounts of these elements for a particular matrix structure are largely a function of section size.

Castability The fluidity of cast iron is a function of its pouring temperature, composition, and eutectic morphology. A higher temperature and higher CE result in better fluidity. Everything else being equal, the fluidity of CG iron is intermediate between that of FG (highest) and SG (lowest) iron (Ref 1). However, because CG iron has a higher strength than FG iron for the same CE, high-CE compositions of CG iron can be used for the pouring of thin castings. Shrinkage Characteristics. With CG irons, obtaining sound castings free from external and internal shrinkage porosity is easier than with SG irons and slightly more difficult than with FG irons. This is because the tendency for mold wall movement also lies between that of SG and FG irons. In relative numbers, solidification expansion has been found to be 4.4 for SG iron and 1 to 1.8 for CG iron if FG iron is 1 (Ref 4). Because of the rather low shrinkage of CG iron, it can sometimes be cast riserless. Expensive pattern changes are therefore not necessary when converting from gray iron to CG iron because the same gating and risering techniques can be applied. Chilling Tendency. Although many believe that the chilling tendency of CG iron is also intermediate between that of FG (lowest) and SG (highest) irons, this is not true. Figure 4 shows the influence of nodularity on the structure of chill pins cast in air set molds (Ref 5). It can be seen that the highest chilling tendency is achieved for irons with nodularities between 6 and 64%. In other words, the chilling tendency of CG iron is higher than that of both SG and FG iron. This correlates with cooling curve data (see the article "Compacted Graphite Irons" in Casting, Volume 15 of ASM Handbook, formerly 9th Edition Metals Handbook) and is explained by the combination of a low nucleation rate and low growth rate occurring during the solidification of CG iron. Fig. 4 Influence of graphite shape over the chilling tendency of cast iron. Type A graphite flake: uniform distribution and random orientation. Type D graphite flake: interdendritic segregation and random orientation (see the article "Classification and Basic Metallurgy of Cast Iron" in this Volume). Source: Ref 5

Mechanical Properties at Room Temperature The in-service behavior of many structural parts is a function not only of their mechanical strength, but also of their deformation properties. Thus it is not surprising to find that many castings fail not because of insufficient strength, but because of a low capacity for deformation. This is especially true under conditions of rapid loading and/or thermal stress. Particularly sensitive to such loading are casting zones that include some defects or abrupt changes in section thickness. The elongation values of about 1% obtainable with high-strength gray iron are insufficient for certain types of applications such as diesel cylinder heads (Ref 6). Compacted graphite irons have strength properties close to those of SG irons, at considerably higher elongations than those of FG iron, and with intermediate thermal conductivities. Consequently, they can successfully outperform other cast irons in a number of applications. The main factors affecting the mechanical properties of CG irons both at room temperatures and at elevated temperatures are: • Composition • Structure (nodularity and matrix) • Section size


Page 104


Compacted Graphite Iron

01 Sep 2005

In turn, the structure is heavily influenced by processing variables such as the type of raw materials, preprocessing of the melt (superheating temperature, holding time, desulfurization), and liquid treatment (graphite compaction and postinoculation). Tensile Properties and Hardness. A comparison between some properties of FG, CG, and SG irons is given in Table 1 . A listing of tensile properties of various CG irons produced by different methods is given in Table 2 . Table 1 Comparison of properties of cerium-treated CG iron with FG iron of the same chemical composition, high-strength pearlitic FG iron, and ferritic SG iron in the as-cast condition High-strength pearlitic FG iron (100% pearlite, 100% FG)(a)

FG iron (100% pearlite, 100% FG)(b)

Chemical composition, %

3.10 C, 2.10 Si, 0.60 Mn

3.61 C, 2.49 Si, 0.05 Mn

36.1 C, 2.54 Si, 0.05 Mn

3.56 C, 2.72 Si, 0.05 Mn

Tensile strength, MPa (ksi)

317 (46)

110 (16)

336 (48.7)

438 (63.5)

0.2% proof stress, MPa (ksi)

...

...

257 (37.3)

285 (41.3)

Elongation, %

...

...

6.7

25.3

108 (15.7)

96.9 (14.05)

158 (22.9)

176 (25.5)

200

156

150

159

Ãat 20 °C (68 °F)

...

...

9.32 (6.87)

24.5 (18.1)

Ãat −20 °C (−4 °F)

...

...

6.57 (4.85)

9.81 (7.23)

Ãat −40 °C (−40 °F)

...

...

7.07 (5.21)

6.18 (4.56)

Ãat 20 °C (68 °F)

4.9

2.0

32.07 (23.7)

176.5 (130.2)

Ãat −20 °C (−4 °F)

...

...

26.48 (19.5)

148.1 (109.2)

Ãat −40 °C (−40 °F)

...

...

26.67 (19.7)

121.6 (89.7)

127.5 (18.5)

49.0 (7.1)

210.8 (30.6)

250.0 (36.3)

Property

Modulus of elasticity, GPA (106 psi) Brinell hardness, HB

Ce-treated CG iron SG iron (100% ferrite, (>95% 80% ferrite, >95% CG)(b) SG, 20% poor SG)(b)

Charpy V-notched-bar impact toughness, J (ft · lbf)

Charpy impact bend toughness, J (ft · lbf)

Rotating-bar fatigue strength, MPa (ksi)

Thermal conductivity, W/(cm · K) 0.419 0.423 0.356 0.327 (a) Mechanical properties determined from a sample with a section size 30 mm (1.2 in.) in diameter. (b) Mechanical properties determined from a Y block 23 mm (0.9 in.) section. Source: Ref 7

Table 2 Tensile properties, hardness, and thermal conductivity of various CG irons at room temperature Tensile strength

0.2% proof stress

Degree of saturation, S C (b)

Graphite type

MPa

ksi

MPa

ksi

As-cast ferrite (>95% F)

1.04

95% CG, 5% SG

336

48.7

257

37.3

Ferritic-pearlitic (>5% P)

1.04

95% CG, 5% SG

298

43.2

224

32.5

As-cast ferrite (90% F, 10% P)

1.00

85% CG, 15% SG

371

53.8

267

38.7

100% ferrite

1.00

85% CG, 15% SG

338

49.0

245

35.5

100% ferrite

1.04−1.09

CG

365 ± 63

53 ± 9

278 ± 42

40 ± 6

...

>90% CG

300−400

43−58

250−300

36−43

1.04

70% CG, 30% SG

320

46.4

242

35

Structural condition(a) Irons treated with additions of cerium

Ferritic-pearlitic (>90% F, 5% P)

5.3

...

...

144

20.9

38.5

128

As-cast ferrite (90% F, 10% P)

5.5

137

19.9

...

...

...

...

100% ferrite

8.0

...

...

...

...

...

140

100% ferrite

7.2 ± 4.5

...

...

...

...

...

138−156

Ferritic-pearlitic (>90% F, 20(f)

>800(f)

>85(f)

>3300(f)

2.0

78


Page 174


Alloy Cast Irons

01 Sep 2005

Low-alloy ferritic

(3.3)

(1.5)

(0.6)

(1.5)

>20(f)

>800(f)

>90(f)

>3500(f)

1.4

54

Low-alloy ferritic

(3.3)

(2.2)

(1.0)

(1.0)

5.8

228

25.9

1020

1.2

47

Low-alloy ferritic

(3.1)

(2.2)

(0.9)

(1.5)

7.2

284

29.0

1140

1.6

62

Type 309 stainless ... ... (25.0) (12.0) nil nil nil nil nil nil (a) Parenthetical values are estimates. Phosphorus and sulfur contents in all iron samples were about 0.10%. (b) Exposure of 2000 h in electric furnace at 760 °C (1400 °F) with air atmosphere containing 17−19% O. (c) Exposed for 492 h in gas-fired heat-treating furnace at 815 °C (1500 °F). (d) 6.05% Cu. (e) 6.0% Cu. (f) Specimen completely burned. Source: Ref 4

Table 13 Oxidation of ferritic and austenitic cast irons and one stainless steel Composition, %(a) Iron

TC

Si

Growth

Cr

Ni

Oxide penetration

mm/yr

mil/yr

mm/yr

mil/yr

After 3723 at 745−760 °C (1375−1400 °F) in electric furnace, air atmosphere Ferritic

3.05

2.67

0.90

1.55

2.0

78

(b)

(b)

Austenitic

2.97

1.63

1.89

20.02

0.8

31

6.9

270

Austenitic

2.52

2.67

5.16

20.03

nil

nil

0.2

6

Austenitic

2.32

1.86

2.86

30.93

nil

nil

2.0

78

Austenitic

1.86

5.84

5.00

29.63

nil

nil

0.12−0.25 incl

0.05

>0.25−0.40 incl

0.06

>0.40−0.55 incl

0.07

>0.55−0.80 incl

0.10

>0.80

0.13

≤0.40

0.15

>0.40−0.50 incl

0.20

>0.50−1.65 incl

0.30

>0.040−0.08 incl

0.03

>0.08−0.13 incl

0.05

>0.050−0.09 incl

0.03

>0.09−0.15 incl

0.05

>0.15−0.23 incl

0.07

>0.23−0.35 incl

0.09

≤0.15

0.08

>0.15−0.20 incl

0.10

>0.20−0.30 incl

0.15

>0.30−0.60 incl

0.20

When copper is required, 0.20% minimum is commonly used

Lead(d) When lead is required, a range of 0.15−0.35 is generally used Note: Boron-treated fine-grain steels are produced to a range of 0.0005−0.003% B. Incl, inclusive.(a) The carbon ranges shown customarily apply when the specified maximum limit for manganese does not exceed 1.10%. When the maximum manganese limit exceeds 1.10%, it is customary to add 0.01 to the carbon range shown. (b) It is not common practice to produce a rephosphorized and resulfurized carbon steel to specified limits for silicon because of its adverse effect on machinability. (c) When silicon is required for rods the following ranges and limits are commonly used: 0.10 max; 0.07−0.15, 0.10−0.20, 0.15−0.35, 0.20−0.40, or 0.30−0.60. (d) Lead is reported only as a range of 0.15−0.35% because it is usually added to the mold or ladle stream as the steel is poured. Source: Ref 1

Table 2 Carbon steel cast or heat chemical limits and ranges Applicable only to structural shapes, plates, strip, sheets, and welded tubing Maximum of Element specified element, % Carbon(a)(b)

Manganese

Phosphorus

Range, %

≤0.15

0.05

>0.15−0.30 incl

0.06

>0.30−0.40 incl

0.07

>0.40−0.60 incl

0.08

>0.60−0.80 incl

0.11

>0.80−1.35 incl

0.14

≤0.50

0.20

>0.050−1.15 incl

0.30

>1.15−1.65 incl

0.35

≤0.08

0.03


Page 235


Sulfur

Silicon

Classification and Designation of Carbon...

01 Sep 2005

>0.08−0.15 incl

0.05

≤0.08

0.03

>0.08−0.15 incl

0.05

>0.15−0.23 incl

0.07

>0.23−0.33 incl

0.10

≤0.15

0.08

>0.15−0.30 incl

0.15

>0.30−0.60 incl

0.30

Copper When copper is required, 0.20% minimum is commonly specified Incl, inclusive.(a) The carbon ranges shown in the range column apply when the specified maximum limit for manganese does not exceed 1.00%. When the maximum manganese limit exceeds 1.00%, add 0.01 to the carbon ranges shown in the table. (b) Maximum of 0.12% C for structural shapes and plates. Source: Ref 1

Table 3 Alloy steel heat composition ranges and limits for bars, blooms, billets, and slabs Range, % Element Carbon

Manganese

Sulfur(a)

Silicon

Chromium

Nickel

Open hearth or basic oxygen steels

Electric furnace steels

≤0.55

0.05

0.05

>0.55−0.70 incl

0.08

0.07

>0.70−0.80 incl

0.10

0.09

>0.80−0.95 incl

0.12

0.11

>0.95−1.35 incl

0.13

0.12

≤0.60

0.20

0.15

>0.60−0.90 incl

0.20

0.20

>0.90−1.05 incl

0.25

0.25

>1.05−1.90 incl

0.30

0.30

>1.90−2.10 incl

0.40

0.35

≤0.050

0.015

0.015

>0.050−0.07 incl

0.02

0.02

>0.07−0.10 incl

0.04

0.04

>0.10−0.14 incl

0.05

0.05

≤0.15

0.08

0.08

>0.15−0.20 incl

0.10

0.10

>0.20−0.40 incl

0.15

0.15

>0.40−0.60 incl

0.20

0.20

>0.60−1.00 incl

0.30

0.30

>1.00−2.20 incl

0.40

0.35

≤0.40

0.15

0.15

>0.40−0.90 incl

0.20

0.20

>0.90−1.05 incl

0.25

0.25

>1.05−1.60 incl

0.30

0.30

>1.60−1.75 incl

(b)

0.35

>1.75−2.10 incl

(b)

0.40

Maximum of specified element, %

>2.10−3.99 incl

(b)

0.50

≤0.50

0.20

0.20

>0.50−1.50 incl

0.30

0.30

>1.50−2.00 incl

0.35

0.35

>2.00−3.00 incl

0.40

0.40


Page 236


Molybdenum

Tungsten

Copper

Vanadium Aluminum


>3.00−5.30 incl

0.50

0.50

>5.30−10.00 incl

1.00

1.00

≤0.10

0.05

0.05

>0.10−0.20 incl

0.07

0.07

>0.20−0.50 incl

0.10

0.10

>0.50−0.80 incl

0.15

0.15

>0.80−1.15 incl

0.20

0.20

≤0.50

0.20

0.20

>0.50−1.00 incl

0.30

0.30

>1.00−2.00 incl

0.50

0.50

>2.00−4.00 incl

0.60

0.60

≤0.60

0.20

0.20

>0.60−1.50 incl

0.30

0.30

>1.50−2.00 incl

0.35

0.35

≤0.25

0.05

0.05

>0.25−0.50 incl

0.10

0.10

≤0.10

0.05

0.05

>0.10−0.20 incl

0.10

0.10

>0.20−0.30 incl

0.15

0.15

>0.30−0.80 incl

0.25

0.25

>0.80−1.30 incl

0.35

0.35

0.45

>1.30−1.80 incl Element Phosphorus

Sulfur

01 Sep 2005

0.45

Steelmaking process

Lowest maximum, %(c)

Basic open hearth, basic oxygen, or basic electric furnace steels

0.035(d)

Basic electric furnace E steels

0.025

Acid open hearth or electric furnace steel

0.050

Basic open hearth, basic oxygen, or basic electric furnace steels

0.040(d)

Basic electric furnace E steels

0.025

Acid open hearth or electric furnace steel 0.050 Inc, inclusive.(a) A range of sulfur content normally indicates a resulfurized steel. (b) Not normally produced by open hearth process. (c) Not applicable to rephosphorized or resulfurized steels. (d) Lower maximum limits on phosphorus and sulfur are required by certain quality descriptors. Source: Ref 2

Table 4 Alloy steel heat composition ranges and limits for plates Range, % Open hearth or basic oxygen steels

Electric furnace steels

≤0.25

0.06

0.05

>0.25−0.40 incl

0.07

0.06

>0.40−0.55 incl

0.08

0.07

>0.55-0.70 incl

0.11

0.10

>0.70

0.14

0.13

Maximum of specified element, %

Element Carbon

Manganese

Sulfur

≤0.45

0.20

0.15

>0.45−0.80 incl

0.25

0.20

>0.80−1.15 incl

0.30

0.25

>1.15−1.70 incl

0.35

0.30

>1.70−2.10 incl

0.40

0.35

≤0.060

0.02

0.02


Page 237


Silicon

Copper

Nickel

Chromium

Molybdenum

Vanadium


01 Sep 2005

>0.060−0.100 incl

0.04

0.04

>0.100−0.140 incl

0.05

0.05

≤0.15

0.08

0.08

>0.15−0.20 incl

0.10

0.10

>0.20−0.40 incl

0.15

0.15

>0.40−0.60 incl

0.20

0.20

>0.60−1.00 incl

0.30

0.30

>1.00−2.20 incl

0.40

0.35

≤0.60

0.20

0.20

>0.60−1.50 incl

0.30

0.30

>1.50−2.00 incl

0.35

0.35

≤0.50

0.20

0.20

>0.50−1.50 incl

0.30

0.30

>1.50−2.00 incl

0.35

0.35

>2.00−3.00 incl

0.40

0.40

>3.00-5.30 incl

0.50

0.50

>5.30−10.00 incl

1.00

1.00

≤0.40

0.20

0.15

>0.40−0.80 incl

0.25

0.20

>0.80−1.05 incl

0.30

0.25

>1.05−1.25 incl

0.35

0.30

>1.25−1.75 incl

0.50

0.40

>1.75−3.99 incl

0.60

0.50

≤0.10

0.05

0.05

>0.10−0.20 incl

0.07

0.07

>0.20−0.50 incl

0.10

0.10

>0.50−0.80 incl

0.15

0.15

0.80−1.15 incl

0.20

0.20

≤0.25

0.05

0.05

0.10 0.10 >0.25−0.50 incl Note: Boron steels can be expected to contain a minimum of 0.0005% B. Alloy steels can be produced with a lead range of 0.15−0.35%. A heat analysis for lead is not determinable because lead is added to the ladle stream while each ingot is poured. Incl, inclusive.Source: Ref 3

Because segregation of some alloying elements is inherent in the solidification of an ingot, different portions will have local chemical compositions that differ slightly from the average composition. Many lengths of bar stock can be made from a single ingot; therefore, some variation in composition between individual bars must be expected. The compositions of individual bars might not conform to the applicable specification, even though the heat analysis does. The chemical composition of an individual bar (or other product) taken from a large heat of steel is called the product analysis or check analysis. Ranges and limits for product analyses are generally broader and less restrictive than the corresponding ranges and limits for heat analyses. Such limits used in standard commercial practice are given in Tables 5 , 6 , and 7 . Table 5 Product analysis tolerances for carbon and alloy steel plates, sheet, piling, and bars for structural applications Tolerance, % Element Carbon Manganese(a)

Under minimum limit

Over maximum limit

≤0.15

0.02

0.03

>0.15−0.40 incl

0.03

0.04

≤0.60

0.05

0.06

>0.60−0.90 incl

0.06

0.08

>0.90−1.20 incl

0.08

0.10

Upper limit or maximum specified value, %


Page 238



01 Sep 2005

>1.20−1.35 incl

0.09

0.11

>1.35−1.65 incl

0.09

0.12

>1.65−1.95 incl

0.11

0.14

>1.95

0.12

0.16

≤0.04

...

0.010

>0.40−0.15 incl

...

(b)

Sulfur

≤0.05

...

0.010

Silicon

≤0.30

0.02

0.03

>0.30−0.40 incl

0.05

0.05

>0.40−2.20 incl

0.06

0.06

≤1.00

0.03

0.03

>1.00−2.00 incl

0.05

0.05

Phosphorus

Nickel Chromium Molybdenum

Copper

≤0.90

0.04

0.04

>0.90−2.10 incl

0.06

0.06

≤0.20

0.01

0.01

>0.20−0.40 incl

0.03

0.03

>0.40−1.15 incl

0.04

0.04

0.20 minimum only

0.02

...

≤1.00

0.03

0.03

>1.00−2.00 incl

0.05

0.05

Titanium

≤0.10

0.01(c)

0.01(c)

Vanadium

≤0.10

0.01(c)

0.01(c)

>0.10−0.25 incl

0.02

0.02

Minimum only specified

0.01

...

Boron

Any

(b)

(b)

Niobium

≤0.10

0.01(c)

0.01(c)

Zirconium

≤0.15

0.03

0.03

Nitrogen 0.005 0.005 ≤0.030 Incl, inclusive.(a) Manganese product analyses tolerances for bars and bar size shapes: ≤0.90, ±0.03; >0.90−2.20 incl, ±0.06. (b) Product analysis not applicable. (c) If the minimum of the range is 0.01%, the under tolerance is 0.005%. Source: Ref 4

Table 6 Product analysis tolerances for carbon and high-strength low-alloy steel bars, blooms, billets, and slabs Tolerance over the maximum limit or under the minimum limit, % Limit or maximum of specified range, %

Element Carbon

Manganese

≤0.065 m2 (100 in.2)

>0.065−0.129 m2 >0.129−0.258 m2 (100−200 in.2) (200−400 in.2) incl incl

>0.258−0.516 m2 (400−800 in.2) incl

≤0.25

0.02

0.03

0.04

0.05

>0.25−0.55 incl

0.03

0.04

0.05

0.06

>0.55

0.04

0.05

0.06

0.07

≤0.90

0.03

0.04

0.06

0.07

>0.90−1.65 incl

0.06

0.06

0.07

0.08

Phosphorus(a)

Over maximum only, ≤0.40

0.008

0.008

0.010

0.015

Sulfur(a)

Over maximum only, ≤0.050

0.008

0.010

0.010

0.015

Silicon

≤0.35

0.02

0.02

0.03

0.04

>0.35−0.60 incl

0.05

...

...

...

Under minimum only

0.02

0.03

...

...

Copper

Lead(b) 0.03 0.03 ... ... 0.15−0.35 incl Note: Rimmed or capped steels and boron are not subject to product analysis tolerances. Product analysis tolerances for alloy elements in


Page 239



01 Sep 2005

high-strength low-alloy steels are given in Table 7 . Incl, inclusive.(a) Because of the degree to which phosphorus and sulfur segregate, product analysis tolerances for those elements are not applicable for rephosphorized and resulfurized steels. (b) Product analysis tolerance for lead applies, both over and under the specified range. Source: Ref 2

Table 7 Product analysis tolerances for alloy steel bars, blooms, billets, and slabs Tolerance over the maximum limit or under the minimum limit for size ranges shown, %

Element Carbon

Manganese

(100 in.2)

>0.065−0.129 m2 (100−200 in.2) incl

>0.129−0.258 m2 (200−400 in.2) incl

>0.258−0.516 m2 (400−800 in.2) incl

≤0.30

0.01

0.02

0.03

0.04

>0.30−0.75 incl

0.02

0.03

0.04

0.05

Limit or maximum of specified range, %

≤0.065 m2

>0.75

0.03

0.04

0.05

0.06

≤0.90

0.03

0.04

0.05

0.06

>0.90−2.10 incl

0.04

0.05

0.06

0.07

Phosphorus

Over max only

0.005

0.010

0.010

0.010

Sulfur

Over max only(a)

0.005

0.010

0.010

0.010

Silicon

≤0.40

0.02

0.02

0.03

0.04

>0.40−2.20 incl

0.05

0.06

0.06

0.07

≤1.00

0.03

0.03

0.03

0.03

>1.00−2.00 incl

0.05

0.05

0.05

0.05

>2.00−5.30 incl

0.07

0.07

0.07

0.07

>5.30−10.00 incl

0.10

0.10

0.10

0.10

≤0.90

0.03

0.04

0.04

0.05

>0.90−2.10 incl

0.05

0.06

0.06

0.07

>2.10−3.99 incl

0.10

0.10

0.12

0.14

≤0.20

0.01

0.01

0.02

0.03

>0.20−0.40 incl

0.02

0.03

0.03

0.04

>0.40−1.15 incl

0.03

0.04

0.05

0.06

≤0.10

0.01

0.01

0.01

0.01

>0.10−0.25 incl

0.02

0.02

0.02

0.02

>0.25−0.50 incl

0.03

0.03

0.03

0.03

Min value specified, check under min limit(b)

0.01

0.01

0.01

0.01

Nickel

Chromium

Molybdenum

Vanadium

≤1.00

0.04

0.05

0.05

0.06

>1.00−4.00 incl

0.08

0.09

0.10

0.12

≤0.10

0.03

...

...

...

>0.10−0.20 incl

0.04

...

...

...

>0.20−0.30 incl

0.05

...

...

...

>0.30−0.80 incl

0.07

...

...

...

>0.80−1.80 incl

0.10

...

...

...

Lead(c)

0.15−0.35 incl

0.03(d)

...

...

...

Copper(c)

≤1.00

0.03

...

...

...

Tungsten Aluminum(c)

0.05

...

...

...

≤0.10

0.01(b)

...

...

...

Niobium(c)

≤0.10

0.01(b)

...

...

...

Zirconium(c)

≤0.15

0.03

...

...

...

>1.00−2.00 incl Titanium(c)

Nitrogen(c) 0.005 ... ... ... ≤0.030 Note: Boron is not subject to product analysis tolerances. Incl, inclusive.(a) Resulfurized steels are not subject to product analysis limits for


Page 240



01 Sep 2005

sulfur. (b) If the minimum of the range is 0.01%, the under tolerance is 0.005%. (c) Tolerances shown apply only to 0.065 m 2 (100 in.2) or less. (d) Tolerance is over and under. Source: Ref 2

Residual elements usually enter steel products from raw materials used to produce pig iron or from scrap steel used in steelmaking. Through careful steelmaking practices, the amounts of these residual elements are generally held to acceptable levels. Sulfur and phosphorus are usually considered deleterious to the mechanical properties of steels; therefore, restrictions are placed on the allowable amounts of these elements for most grades. The amounts of sulfur and phosphorus are invariably reported in the analyses of both carbon and alloy steels. Other residual alloying elements generally exert a lesser influence than sulfur and phosphorus on the properties of steel. For many grades of steel, limitations on the amounts of these residual elements are either optional or omitted entirely. Amounts of residual alloying elements are generally not reported in either heat or product analyses, except for special reasons. Silicon Content of Steels. The composition requirements for many steels, particularly plain carbon steels, contain no specific restriction on silicon content. The lack of a silicon requirement is not an omission, but instead indicates recognition that the amount of silicon in a steel can often be traced directly to the deoxidation practice employed in making it (further information can be found in the section "Types of Steel Based on Deoxidation Practice" in this article). Rimmed and capped steels are not deoxidized; the only silicon present is the residual amount left from scrap or raw materials, typically less than 0.05% Si. Specifications and orders for these steels customarily indicate that the steel must be made rimmed or capped, as required by the purchaser, restrictions on silicon content are not usually given. The extent of rimming action during the solidification of semikilled steel ingots must be carefully controlled by matching the amount of deoxidizer with the oxygen content of the molten steel. The amount of silicon required for deoxidation may vary from heat to heat. Thus, the silicon content of the solid metal can also vary slightly from heat to heat. A maximum silicon content of 0.10% is sometimes specified for semikilled steel, but this requirement is not very restrictive; for certain heats, a silicon addition sufficient to leave a residue of 0.10% may be enough of an addition to kill the steel. Killed steels are fully deoxidized during their manufacture; deoxidation can be accomplished by additions of silicon, aluminum, or both, or by vacuum treatment of the molten steel. Because it is the least costly of these methods, silicon deoxidation is frequently used. For silicon-killed steels, a range of 0.15 to 0.30% Si is often specified, providing the manufacturer with adequate flexibility to compensate for variations in the steelmaking process and ensuring a steel acceptable for most applications. Aluminum-killed or vacuum-deoxidized steels require no silicon; a requirement for minimum silicon content in such steel is unnecessary. A maximum permissible silicon content is appropriate for all killed plain carbon steels; a minimum silicon content implies a restriction that the steel must be silicon killed. Silicon is intentionally added to some alloy steels, for which it serves as both a deoxidizer and an alloying element to modify the properties of the steel. An acceptable range of silicon content would be appropriate for these steels. Users and specifiers of steel mill products must realize that the silicon content of these items cannot be established independently of deoxidation practice. In ordering mill products, it is often desirable to cite a standard specification (such as an ASTM specification) where the various ramifications of restrictions on silicon content have already been considered in preparing the specification. In some instances, such as the forming of low-carbon steel sheet, the choice of deoxidation practice can significantly affect the performance of the steel; in such cases, it is appropriate to specify the desired practice. Types of Steel Based on Deoxidation Practice (Ref 3) Steels, when cast into ingots, can be classified into four types based on the deoxidation practice employed or, alternatively, by the amount of gas evolved during solidification. These types are killed, semikilled, rimmed, or capped steels (Fig. 2 ). Fig. 2 Eight typical conditions of commercial steel ingots, cast in identical bottle-top molds, in relation to the degree of suppression of gas evolution. The dotted line indicates the height to which the steel originally was poured in each ingot mold. Depending on the carbon and, more importantly, the oxygen content of the steel, the ingot structures range from that of a fully killed ingot (No. 1) to that of a violently rimmed ingot (No. 8). Source:Ref 5

Killed steel is a type of steel from which there is only a slight evolution of gases during solidification of the metal after pouring. Killed steels are characterized by more uniform chemical composition and properties as compared to the other types. Alloy steels, forging steels, and steels for carburizing are generally killed.


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01 Sep 2005

Killed steel is produced by various steel-melting practices involving the use of certain deoxidizing elements which act with varying intensities. The most common of these are silicon and aluminum; however, vanadium, titanium, and zirconium are sometimes used. Deoxidation practices in the manufacture of killed steels are normally left to the discretion of the producer. Semikilled steel is a type of steel wherein there is a greater degree of gas evolution than in killed steel but less than in capped or rimmed steel. The amount of deoxidizer used (customarily silicon or aluminum) will determine the amount of gas evolved. Semikilled steels generally have a carbon content within the range of 0.15 to 0.30%; they are used for a wide range of structural shape applications. Semikilled steels are characterized by variable degrees of uniformity in composition, which are intermediate between those of killed and rimmed steels. Semikilled steel has a pronounced tendency for positive chemical segregation at the top-center of the ingot (Fig. 2 ). Rimmed Steels. In the production of rimmed steels, no deoxidizing agents are added in the furnace. These steels are characterized by marked differences in chemical composition across the section and from the top to the bottom of the ingot (Fig. 2 ). They have an outer rim that is lower in carbon, phosphorus, and sulfur than the average composition of the whole ingot, and an inner portion, or core, that has higher levels than the average of those elements. The typical structure of the rimmed steel ingot results from a marked gas evolution during solidification of the outer rim. During the solidification of the rim, the concentration of certain elements increases in the liquid portion of the ingot. During solidification of the core, some increase in segregation occurs in the upper and central portions of the ingot. The structural pattern of the ingot persists through the rolling process to the final product (rimmed ingots are best suited for steel sheets). The technology of manufacturing rimmed steels limits the maximum content of carbon and manganese, and those maximums vary among producers. Rimmed steels do not retain any significant percentages of highly oxidizable elements such as aluminum, silicon, or titanium. Capped steels have characteristics similar to those of rimmed steels but to a degree intermediate between those of rimmed and semikilled steels. A deoxidizer may be added to effect a controlled rimming action when the ingot is cast. The gas entrapped during solidification is in excess of that needed to counteract normal shrinkage, resulting in a tendency for the steel to rise in the mold. The capping operation limits the time of gas evolution and prevents the formation of an excessive number of gas voids within the ingot. Mechanically capped steel is cast in bottle-top molds using a heavy metal cap. Chemically capped steel is cast in open-top molds. The capping is accomplished by adding aluminum or ferrosilicon to the top of the ingot, causing the steel at the top surface to solidify rapidly. The top portion of the ingot is discarded. The capped ingot practice is usually applied to steel with carbon contents greater than 0.15% that is used for sheet, strip, wire, and bars. Quality Descriptors The need for communication among producers and between producers and users has resulted in the development of a group of terms known as fundamental quality descriptors. These are names applied to various steel products to imply that the particular products possess certain characteristics that make them especially well suited for specific applications or fabrication processes. The fundamental quality descriptors in common use are listed in Table 8 . Table 8 Quality descriptions of carbon and alloy steels Carbon steels Semifinished for forging Forging quality ÃSpecial hardenability ÃSpecial internal soundness ÃNonmetallic inclusion requirement ÃSpecial surface Carbon steel structural sections Structural quality Carbon steel plates Regular quality Structural quality Cold-drawing quality Cold-pressing quality Cold-flanging quality Forging quality Pressure vessel quality Hot-rolled carbon steel bars


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01 Sep 2005

Merchant quality Special quality ÃSpecial hardenability ÃSpecial internal soundness ÃNonmetallic inclusion requirement ÃSpecial surface Scrapless nut quality Axle shaft quality Cold extrusion quality Cold-heading and cold-forging quality Cold-finished carbon steel bars Standard quality ÃSpecial hardenability ÃSpecial internal soundness ÃNonmetallic inclusion requirement ÃSpecial surface Cold-heading and cold-forging quality Cold extrusion quality Hot-rolled sheets Commercial quality Drawing quality Drawing quality special killed Structural quality Col-rolled sheets Commercial quality Drawing quality Drawing quality special killed Structural quality Porcelain enameling sheets Commercial quality Drawing quality Drawing quality special killed Long terne sheets Commercial quality Drawing quality Drawing quality special killed Structural quality Galvanized sheets Commercial quality Drawing quality Drawing quality special killed Lock-forming quality Electrolytic zinc coated sheets Commercial quality Drawing quality Drawing quality special killed Structural quality Hot-rolled strip Commercial quality


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01 Sep 2005

Drawing quality Drawing quality special killed Structural quality Cold-rolled strip Specific quality descriptions are not provided in cold-rolled strip because this product is largely produced for specific end use Tin mill products Specific quality descriptions are not applicable to tin mill products Carbon steel wire Industrial quality wire Cold extrusion wires Heading, forging, and roll-threading wires Mechanical spring wires Upholstery spring construction wires Welding wire Carbon steel flat wire Stitching wire Stapling wire Carbon steel pipe Structural tubing Line pipe Oil country tubular goods Steel specialty tubular products Pressure tubing Mechanical tubing Aircraft tubing Hot-rolled carbon steel wire rods Industrial quality Rods for manufacture of wire intended for electric welded chain Rods for heading, forging, and roll-threading wire Rods for lock washer wire Rods for scrapless nut wire Rods for upholstery spring wire Rods for welding wire Alloy steels Alloy steel plates Drawing quality Pressure vessel quality Structural quality Aircraft physical quality Hot-rolled alloy steel bars Regular quality Aircraft quality or steel subject to magnetic particle inspection Axle shaft quality Bearing quality Cold-heading quality Special cold-heading quality Rifle barrel quality, gun quality, shell or A.P. shot quality Alloy steel wire Aircraft quality


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01 Sep 2005

Bearing quality Special surface quality Cold-finished alloy steel bars Regular quality Aircraft quality or steel subject to magnetic particle inspection Axle shaft quality Bearing shaft quality Cold-heading quality Special cold-heading quality Rifle barrel quality, gun quality, shell or A.P. shot quality Line pipe Oil country tubular goods Steel specialty tubular goods Pressure tubing Mechanical tubing Stainless and head-resisting pipe, pressure tubing, and mechanical tubing Aircraft tubing Pipe Source: Ref 6

Some of the quality descriptors listed in Table 8 such as forging quality or cold extrusion quality are self-explanatory. The meaning of others is less obvious: for example, merchant quality hot-rolled carbon steel bars are made for noncritical applications requiring modest strength and mild bending or forming, but not requiring forging or heat treating. The descriptor for one particular steel commodity is not necessarily carried over to subsequent products made from that commodity⎯for example, standard quality cold-finished bars are made from special quality hot-rolled bars. The various mechanical and physical attributes implied by a quality descriptor arise from the combined effects of several factors, including: • • • • • • •

The degree of internal soundess The relative uniformity of chemical composition The relative freedom from surface imperfections The size of the discard cropped from the ingot Extensive testing during manufacture The number, size, and distribution of nonmetallic inclusions Hardenability requirements

Control of these factors during manufacture is necessary to achieve mill products having the desired characteristics. The extent of the control over these and other related factors is another piece of information conveyed by the quality descriptor. Some, but not all, of the fundamental descriptors may be modified by one or more additional requirements, as may be appropriate: special discard, macroetch test, restricted chemical composition, maximum incidental (residual) alloy, special hardenability or austenitic grain size. These restrictions could be applied to forging quality alloy steel bars, but not to merchant quality bars. Understanding the various quality descriptors is complicated by the fact that most of the requirements that qualify a steel for a particular descriptor are subjective. Only nonmetallic inclusion count, restrictions on chemical composition ranges and incidental alloying elements, austenitic grain size, and special hardenability are quantified. The subjective evaluation of the other characteristics depends on the skill and experience of those who make the evaluation. Although the use of these subjective quality descriptors might seem imprecise and unworkable, steel products made to meet the requirements of a particular quality descriptor can be relied upon to have those characteristics necessary for that product to be used in the indicated application or fabrication operation.

Effects of Alloying Elements (Ref 6) Steels form one of the most complex group of alloys in common use. The synergistic effect of alloying elements and heat treatment produce a tremendous variety of microstructures and properties (characteristics). Given the limited scope of this article, it would be impossible to include a detailed survey of the effects of alloying elements on the iron-carbon equilibrium diagram. This complicated subject, which is briefly reviewed in the article "Microstructures, Processing, and Properties of Steels" in this Volume, lies in the domain of ferrous physical metallurgy and has also been reviewed extensively in the literature (Ref 7, 8, 9,


Page 245



01 Sep 2005

10, 11). In this section, the effects of various elements on steelmaking (deoxidation) practices and steel characteristics will be briefly outlined. It should be noted that the effects of a single alloying elements are modified by the influence of other elements. These interrelations must be considered when evaluating a change in the composition of a steel. For the sake of simplicity, however, the various alloying elements listed below are discussed separately. Carbon. The amount of carbon required in the finished steel limits the type of steel that can be made. As the carbon content of rimmed steels increases, surface quality becomes impaired. Killed steels in approximately the 0.15 to 0.30% C content level may have poorer surface quality and require special processing to attain surface quality comparable to steels with higher or lower carbon contents. Carbon has a moderate tendency to segregate, and carbon segregation is often more significant than the segregation of other elements. Carbon, which has a major effect on steel properties, is the principal hardening element in all steel. Tensile strength in the as-rolled condition increases as carbon content increases (up to about 0.85% C). Ductility and weldability decrease with increasing carbon. Manganese has less of a tendency toward macrosegregation than any of the common elements. Steels above 0.60% Mn cannot be readily rimmed. Manganese is beneficial to surface quality in all carbon ranges (with the exception of extremely low carbon rimmed steels) and is particularly beneficial in resulfurized steels. It contributes to strength and hardness, but to a lesser degree than does carbon; the amount of increase is dependent upon the carbon content. Increasing the manganese content decreases ductility and weldability, but to a lesser extent than does carbon. Manganese has a strong effect on increasing the hardenability of a steel. Phosphorus segregates, but to a lesser degree than carbon and sulfur. Increasing phosphorus increases strength and hardness and decreases ductility and notch impact toughness in the as-rolled condition. The decreases in ductility and toughness are greater in quenched and tempered higher-carbon steels. Higher phosphorus is often specified in low-carbon free-machining steels to improve machinability (see the article "Machinability of Steels" in this Volume). Sulfur. Increased sulfur content lowers transverse ductility and notch impact toughness but has only a slight effect on longitudinal mechanical properties. Weldability decreases with increasing sulfur content. This element is very detrimental to surface quality, particularly in the lower-carbon and lower-manganese steels. For these reasons, only a maximum limit is specified for most steels. The only exception is the group of free-machining steels, where sulfur is added to improve machinability; in this case a range is specified (see the article "Machinability of Steels" in this Volume). Sulfur has a greater segregation tendency than any of the other common elements. Sulfur occurs in steel principally in the form of sulfide inclusions. Obviously, a greater frequency of such inclusions can be expected in the resulfurized grades. Silicon is one of the principal deoxidizers used in steelmaking; therefore, the amount of silicon present is related to the type of steel. Rimmed and capped steels contain no significant amounts of silicon. Semikilled steels may contain moderate amounts of silicon, although there is a definite maximum amount that can be tolerated in such steels. Killed carbon steels may contain any amount of silicon up to 0.60% maximum. Silicon is somewhat less effective than manganese in increasing as-rolled strength and hardness. Silicon has only a slight tendency to segregate. In low-carbon steels, silicon is usually detrimental to surface quality, and this condition is more pronounced in low-carbon resulfurized grades. Copper has a moderate tendency to segregate. Copper in appreciable amounts is detrimental to hot-working operations. Copper adversely affects forge welding, but it does not seriously affect arc or oxyacetylene welding. Copper is detrimental to surface quality and exaggerates the surface defects inherent in resulfurized steels. Copper is, however, beneficial to atmospheric corrosion resistance when present in amounts exceeding 0.20%. Steels containing these levels of copper are referred to as weathering steels and are described in the article "High-Strength Structural and High-Strength Low-Alloy Steels" in this Volume; they are also included in the descriptions of high-strength low-alloy steels given later in this article. Lead is sometimes added to carbon and alloy steels through mechanical dispersion during teeming for the purpose of improving the machining characteristics of the steels. These additions are generally in the range of 0.15 to 0.35% (see the article "Machinability of Steels" in this Volume for details). Boron is added to fully killed steel to improve hardenability. Boron-treated steels are produced to a range of 0.0005 to 0.003%. Whenever boron is substituted in part for other alloys, it should be done only with hardenability in mind because the lowered alloy content may be harmful for some applications. Boron is most effective in lower carbon steels. Boron steels are discussed in the Section "Hardenability of Carbon and Low-Alloy Steels" in this Volume. Chromium is generally added to steel to increase resistance to corrosion and oxidation, to increase hardenability, to improve high-temperature strength, or to improve abrasion resistance in high-carbon compositions. Chromium is a strong carbide former. Complex chromium-iron carbides go into solution in austenite slowly; therefore, a sufficient heating time before quenching is necessary. Chromium can be used as a hardening element, and is frequently used with a toughening element such as nickel to producesuperior mechanical properties. At higher temperatures, chromium contributes increased strength; it is ordinarily used for applications of this nature in conjunction with molybdenum. Nickel, when used as an alloying element in constructional steels, is a ferrite strengthener. Because nickel does not form any carbide compounds in steel, it remains in solution in the ferrite, thus strengthening and toughening the ferrite phase. Nickel steels are easily heat treated because nickel lowers the critical cooling rate. In combination with chromium, nickel produces alloy steels with greater hardenability, higher impact strength, and greater fatigue resistance than can be achieved in carbon steels. Molybdenum is added to constructional steels in the normal amounts of 0.10 to 1.00%. When molybdenum is in solid solution in austenite prior to quenching, the reaction rates for transformation become considerably slower as compared with carbon steel. Molybdenum can induce secondary hardening during the tempering of quenched steels and enhances the creep


Page 246



01 Sep 2005

strength of low-alloy steels at elevated temperatures. Alloy steels that contain 0.15 to 0.30% Mo display a minimized susceptibility to temper embrittlement (see the article "Embrittlement of Steels" in this Volume for a discussion of temper embrittlement and other forms of thermal embrittlement). Niobium. Small additions of niobium increase the yield strength and, to a lesser degree, the tensile strength of carbon steel. The addition of 0.02% Nb can increase the yield strength of medium-carbon steel by 70 to 100 MPa (10 to 15 ksi). This increased strength may be accompanied by considerably impaired notch toughness unless special measures are used to refine grain size during hot rolling. Grain refinement during hot rolling involves special thermomechanical processing techniques such as controlled rolling practices, low finishing temperatures for final reduction passes, and accelerated cooling after rolling is completed (further discussion of controlled rolling can be found in the article "High-Strength Structural and High-Strength Low-Alloy Steels" in this Volume). Aluminum is widely used as a deoxidizer and for control of grain size. When added to steel in specified amounts, it controls austenite grain growth in reheated steels. Of all the alloying elements, aluminum is the most effective in controlling grain growth prior to quenching. Titanium, zirconium, and vanadium are also effective grain growth inhibitors; however, for structural grades that are heat treated (quenched and tempered), these three elements may have adverse effects on hardenability because their carbides are quite stable and difficult to dissolve in austenite prior to quenching. Titanium and Zirconium. The effects of titanium are similar to those of vanadium and niobium, but it is only useful in fully killed (aluminum-deoxidized) steels because of its strong deoxidizing effects. Zirconium can also be added to killed high-strength low-alloy steels to obtain improvements in inclusion characteristics, particularly sulfide inclusions where changes in inclusion shape improve ductility in transverse bending.

Carbon Steels The American Iron and Steel Institute defines carbon steel as follows (Ref 2, 3): Steel is considered to be carbon steel when no minimum content is specified or required for chromium, cobalt, columbium [niobium], molybdenum, nickel, titanium, tungsten, vanadium or zirconium, or any other element to be added to obtain a desired alloying effect; when the specified minimum for copper does not exceed 0.40 per cent; or when the maximum content specified for any of the following elements does not exceed the percentages noted: manganese 1.65, silicon 0.60, copper 0.60.

Carbon steel can be classified, according to various deoxidation practices, as rimmed, capped, semikilled, or killed steel. Deoxidation practice and the steelmaking process will have an effect on the characteristics and properties of the steel (see the article "Steel Processing Technology" in this Volume). However, variations in carbon have the greatest effect on mechanical properties, with increasing carbon content leading to increase hardness and strength (see the article "Microstructures, Processing, and Properties of Steels" in this Volume). As such, carbon steels are generally categorized according to their carbon content. Generally speaking, carbon steels contain up to 2% total alloying elements and can be subdivided into low-carbon steels, medium-carbon steels, high-carbon steels, and ultrahigh-carbon steels; each of these designations is discussed below. As a group, carbon steels are by far the most frequently used steel. Tables 9 and 10 indicate that more than 85% of the steel produced and shipped in the United States is carbon steel. Chemical compositions for carbon steels are provided in the tables referenced in the section "SAE-AISI Designations" in this article. See Tables 11 , 12 , 13 , 14 , 15 , 16 , 17 , 18 , 19 , 20 , 21 , and 22 . Table 9 Raw steel production by type of furnace, grade, and cast Total production Total all grades, net tons × 103 Carbo n Year

By grade, %

By type of furnace, %

Production by type of cast, net tons × 103 Continuo us Steel Ingots castings castings

Alloy

Stainles s

Open heart

Basic oxygen process

Electric

10.9

2.2

5.1

58.0

36.9

38,615

61,232

77

87.5

10.2

2.3

3.0

58.9

38.1

35,802

53,284

65

87.5

10.4

2.1

4.1

58.7

37.2

36,487

45,064

55

88,259

86.9

11.2

1.9

7.3

58.8

33.9

49,035

39,161

63

92,528

86.4

11.7

1.9

9.0

57.1

33.9

55,787

36,669

74

Alloy

Stainl ess

Total

Carbon

1988 86,823 10,902

2199

99,924

86.9

1987 77,976

9,147

2028

89,151

1986 71,413

8,505

1689

81,606

1985 76,699

9,877

1683

1984 79,918 10,838 Source: Ref 12

1772

Table 10 Net shipments of United States steel mill products, all grades 1988 Steel products Ingots and steel for castings Blooms, slabs, and billets Skelp


1987

Net tons × 103

%

Net tons × 103

%

385

0.5

381

0.5

1,542

1.8

1,212

1.6

(a)

...

22

...

Page 247



01 Sep 2005

Wire rods

4,048

4.8

3,840

5.0

Structural shapes (≥75 mm, or 3 in.)

4,860

5.8

4,839

6.3

Steel piling

349

0.4

280

0.4

Plates cut in lengths

5,044

6.0

4,048

5.3

Plates in coils

2,284

2.7

(b)

...

Ãstandard (>27 kg, or 60 lb)

460

0.5

351

0.5

Ãall other

37

0.0

15

...

Railroad accessories

118

0.1

62

0.1

Wheels (rolled and forged)

(a)

...

58

0.1

Axles

(a)

...

29

...

6,460

7.7

6,048

7.9

Rails

Bars Ãhot rolled Ãbar-size light shapes

1,373

1.6

1,190

1.6

Ãreinforcing

5,091

6.1

4,918

6.4

Ãcold finished

1,499

1.8

1,361

1.8

64

0.1

58

0.1

Ãstandard

1,238

1.5

969

1.3

Ãoil country goods

1,130

1.3

919

1.2

Ãline

808

1.0

620

0.8

Ãmechanical

901

1.1

767

1.0

Ãpressure

59

0.1

72

0.1

Ãstructural

178

0.2

180

0.2

Tool steel Pipe and tubing

Ãpipe for piling

74

0.1

(c)

...

Ãstainless

55

0.1

42

0.1

1,073

1.3

800

1.0

Ãnails and staples

(a)

...

218

0.3

Ãbarbed and twisted

(a)

...

49

0.1

woven wire fence

(a)

...

13

...

Ãbale ties and baling wire

(a)

...

25

...

Black plate

283

0.3

205

0.3

Wire Ãdrawn

Tin plate

2,806

3.3

2,765

3.6

Tin free steel

899

1.1

939

1.2

Tin coated sheets

81

0.1

79

0.1

Ãhot rolled

12,589

15.0

13,048

17.0

Ãcold rolled

13,871

16.5

13,859

18.1

Ãgalvanized, hot dipped

8,115

9.7

7,660

10.0

Ãgalvanized, electrolytic

2,134

2.5

1,432

1.9

Ãall other metallic coated

1,262

1.5

1,228

1.6

524

0.6

465

0.6

1,203

1.4

657

0.9

Sheets

Sheets and strip

Ãelectrical Strip Ãhot rolled Ãcold rolled

941

1.1

929

1.2

Total steel mill products

83,840

100.0

76,654

100.0

Carbon

77,702

92.7

68,116

88.9


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Stainless and heat resisting

1,586

1.9

01 Sep 2005

1,418

1.8

Alloy (other than stainless) 4,552 5.4 7,120 9.3 Source: Ref 12(a) Effective 1 January 1988, these products are no longer classified as steel mill products by AISI. Consequently, comparable shipment tonnage is now included in applicable semifinished forms or drawn wire. (b) Prior to 1988 included in sheets hot rolled. (c) Prior to 1988 included in structural pipe and tubing. Source: Ref 12

Table 11 SAE-AISI system of designations

Table 12 Carbon steel compositions Applicable to semifinished products for forging, hot-rolled and cold-finished bars, wire rods, and seamless tubing Cast or heat chemical ranges and limits, %(a) UNS SAE-AISI number

number

C

Mn

P max

S max

G10050

1005

0.06 max

0.35 max

0.040

0.050

G10060

1006

0.08 max

0.25−0.40

0.040

0.50

G10080

1008

0.10 max

0.30−0.50

0.040

0.050

G10100

1010

0.08−0.13

0.30−0.60

0.040

0.050

G10120

1012

0.10−0.15

0.30−0.60

0.040

0.050

G10130

1013

0.11−0.16

0.50−0.80

0.040

0.050

G10150

1015

0.13−018

0.30−0.60

0.040

0.050

G10160

1016

013−0.18

0.60−0.90

0.040

0.050

G10170

1017

0.15−0.20

0.30−0.60

0.040

0.050

G10180

1018

0.15−0.20

0.60−0.90

0.040

0.050

G10190

1019

0.15−0.20

0.70−1.00

0.040

0.050

G10200

1020

0.18−0.23

0.30−0.60

0.040

0.050

G10210

1021

0.18−0.23

0.60−0.90

0.040

0.050

G10220

1022

0.18−0.23

0.70−1.00

0.040

0.050

G10230

1023

0.20−0.25

0.30−0.60

0.040

0.050

G10250

1025

0.22−0.28

0.30−0.60

0.040

0.050

G10260

1026

0.22−0.28

0.60−0.90

0.040

0.050

G10290

1029

0.25−0.31

0.60−0.90

0.040

0.050

G10300

1030

0.28−0.34

0.60−0.90

0.040

0.050


Page 249



01 Sep 2005

G10350

1035

0.32−0.38

0.60−0.90

0.040

0.050

G10370

1037

0.32−0.38

0.70−1.00

0.040

0.050

G10380

1038

0.35−0.42

0.60−0.90

0.040

0.050

G10390

1039

0.37−0.44

0.70−1.00

0.040

0.050

G10400

1040

0.37−0.44

0.60−0.90

0.040

0.050

G10420

1042

0.40−0.47

0.60−0.90

0.040

0.050

G10430

1043

0.40−0.47

0.70−1.00

0.040

0.050

G10440

1044

0.43−0.50

0.30−0.60

0.040

0.050

G10450

1045

0.43−0.50

0.60−0.90

0.040

0.050

G10460

1046

0.43−0.50

0.70−1.00

0.040

0.050

G10490

1049

0.46−0.53

0.60−0.90

0.040

0.050

G10500

1050

0.48−0.55

0.60−0.90

0.040

0.050

G10530

1053

0.48−0.55

0.70−1.00

0.040

0.050

G10550

1055

0.50−0.60

0.60−0.90

0.040

0.050

G10590

1059

0.55−0.65

0.50−0.80

0.040

0.050

G10600

1060

0.55−0.65

0.60−0.90

0.040

0.050

G10640

1064

0.60−0.70

0.50−0.80

0.040

0.050

G10650

1065

0.60−0.70

0.60−0.90

0.040

0.050

G10690

1069

0.65−0.75

0.40−0.70

0.040

0.050

G10700

1070

0.65−0.75

0.60−0.90

0.040

0.050

G10740

1074

0.70−0.80

0.50−0.80

0.040

0.050

G10750

1075

0.70−0.80

0.40−0.70

0.040

0.050

G10780

1078

0.72−0.85

0.30−0.60

0.040

0.050

G10800

1080

0.75−0.88

0.60−0.90

0.040

0.050

G10840

1084

0.80−0.93

0.60−0.90

0.040

0.050

G10850

1085

0.80−0.93

0.70−1.00

0.040

0.050

G10860

1086

0.80−0.93

0.30−0.50

0.040

0.050

G10900

1090

0.85−0.98

0.60−0.90

0.040

0.050

G10950 1095 0.040 0.050 0.90−1.03 0.30−0.50 (a) When silicon ranges or limits are required for bar and semifinished products, the values in Table 1 apply. For rods, the following ranges are commonly used: 0.10 max; 0.07−0.15%; 0.10−0.20%; 0.15−0.35%; 0.20−0.40%; and 0.30−0.60%. Steels listed in this table can be produced with additions of lead or boron. Leaded steels typically contain 0.15−0.35% Pb and are identified by inserting the letter L in the designation (10L45); boron steels can be expected to contain 0.0005−0.003% B and are identified by inserting the letter B in the designation (10B46). Source: Ref 1

Table 13 Carbon steel compositions Applicable only to structural shapes, plates, strip, sheets, and welded tubing Cast or heat chemical ranges and limits, %(a) UNS SAE-AISI number

number

C

Mn

P max

S max

G10060

1006

0.08 max

0.45 max

0.040

0.050

G10080

1008

0.10 max

0.50 max

0.040

0.050

G10090

1009

0.15 max

0.60 max

0.040

0.050

G10100

1010

0.08−0.13

0.30−0.60

0.040

0.050

G10120

1012

0.10−0.15

0.30−0.60

0.040

0.050

G10150

1015

0.12−0.18

0.30−0.60

0.040

0.050

G10160

1016

0.12−0.18

0.60−0.90

0.040

0.050

G10170

1017

0.14−0.20

0.30−0.60

0.040

0.050

G10180

1018

0.60−0.90

0.040

0.050


0.14−0.20

Page 250



01 Sep 2005

G10190

1019

0.14−0.20

0.70−1.00

0.040

0.050

G10200

1020

0.17−0.23

0.30−0.60

0.040

0.050

G10210

1021

0.17−0.23

0.60−0.90

0.040

0.050

G10220

1022

0.17−0.23

0.70−1.00

0.040

0.050

G10230

1023

0.19−0.25

0.30−0.60

0.040

0.050

G10250

1025

0.22−0.28

0.30−0.60

0.040

0.050

G10260

1026

0.22−0.28

0.60−0.90

0.040

0.050

G10300

1030

0.27−0.34

0.60−0.90

0.040

0.050

G10330

1033

0.29−0.36

0.70−1.00

0.040

0.050

G10350

1035

0.31−0.38

0.60−0.90

0.040

0.050

G10370

1037

0.31−0.38

0.70−1.00

0.040

0.050

G10380

1038

0.34−0.42

0.60−0.90

0.040

0.050

G10390

1039

0.36−0.44

0.70−1.00

0.040

0.050

G10400

1040

0.36−0.44

0.60−0.90

0.040

0.050

G10420

1042

0.39−0.47

0.60−0.90

0.040

0.050

G10430

1043

0.39−0.47

0.70−1.00

0.040

0.050

G10450

1045

0.42−0.50

0.60−0.90

0.040

0.050

G10460

1046

0.42−0.50

0.70−1.00

0.040

0.050

G10490

1049

0.45−0.53

0.60−0.90

0.040

0.050

G10500

1050

0.47−0.55

0.60−0.90

0.040

0.050

G10550

1055

0.52−0.60

0.60−0.90

0.040

0.050

G10600

1060

0.55−0.66

0.60−0.90

0.040

0.050

G10640

1064

0.59−0.70

0.50−0.80

0.040

0.050

G10650

1065

0.59−0.70

0.60−0.90

0.040

0.050

G10700

1070

0.65−0.76

0.60−0.90

0.040

0.050

G10740

1074

0.69−0.80

0.50−0.80

0.040

0.050

G10750

1075

0.69−0.80

0.40−0.70

0.040

0.050

G10780

1078

0.72−0.86

0.30−0.60

0.040

0.050

G10800

1080

0.74−0.88

0.60−0.90

0.040

0.050

G10840

1084

0.80−0.94

0.60−0.90

0.040

0.050

G10850

1085

0.80−0.94

0.70−1.00

0.040

0.050

G10860

1086

0.80−0.94

0.30−0.50

0.040

0.050

G10900

1090

0.60−0.90

0.040

0.050

0.84−0.98

G10950 1095 0.040 0.050 0.90−1.04 0.30−0.50 (a) When silicon ranges or limits are required, the following ranges and limits are commonly used: up to SAE 1025 inclusive, 0.10% max, 0.10−0.25%, or 0.15−0.35%. Over SAE 1025, 0.10−0.25% or 0.15−0.35%. Source: Ref 1

Table 14 Composition ranges and limits for merchant quality steels SAE-AISI number M1008

Cast or heat chemical ranges and limits, % (a) C

Mn

P max

S max

0.10 max

0.25−0.60

0.04

0.05 0.05

M1010

0.07−0.14

0.25−0.60

0.04

M1012

0.09−0.16

0.25−0.60

0.04

0.05

M1015

0.12−0.19

0.25−0.60

0.04

0.05

M1017

0.14−0.21

0.25−0.60

0.04

0.05

M1020

0.17−0.24

0.25−0.60

0.04

0.05


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01 Sep 2005

M1023

0.19−0.27

0.25−0.60

0.04

0.05

M1025

0.20−0.30

0.25−0.60

0.04

0.05

M1031

0.26−0.36

0.25−0.60

0.04

0.05

0.04

0.05

M1044 0.40−0.50 0.25−0.60 (a) Merchant quality steel bars are not produced to any specified silicon content. Source: Ref 1

Table 15 Free-cutting (resulfurized) carbon steel compositions Applicable to semifinished products for forging, hot-rolled and cold-finished bars, wire rods, and seamless tubing Cast or heat chemical ranges and limits, %(a) UNS SAE-AISI number

number

C

Mn

P max

S

G11080

1108

0.08−0.13

0.50−0.80

0.040

0.08−0.13

G11100

1110

0.08−0.13

0.30−0.60

0.040

0.08−0.13

G11170

1117

0.14−0.20

1.00−1.30

0.040

0.08−0.13

G11180

1118

0.14−0.20

1.30−1.60

0.040

0.08−0.13

G11370

1137

0.32−0.39

1.35−1.65

0.040

0.08−0.13

G11390

1139

0.35−0.43

1.35−1.65

0.040

0.13−0.20

G11400

1140

0.37−0.44

0.70−1.00

0.040

0.08−0.13

G11410

1141

0.37−0.45

1.35−1.65

0.040

0.08−0.13

G11440

1144

0.40−0.48

1.35−1.65

0.040

0.24−0.33

G11460

1146

0.42−0.49

0.70−1.00

0.040

0.08−0.13

G11510 1151 0.040 0.48−0.55 0.70−1.00 0.08−0.13 (a) When lead ranges or limits are required, or when silicon ranges or limits are required for bars or semifinished products, the values in Table 1 apply. For rods, the following ranges and limits for silicon are commonly used: up to SAE 1110 inclusive, 0.10% max; SAE 1117 and over, 0.10% max, 0.10−0.20%, or 0.15−0.35%. Source: Ref 1

Table 16 Free-cutting (rephosphorized and resulfurized) carbon steel compositions Applicable to semifinished products for forging, hot-rolled and cold-finished bars, wire rods, and seamless tubing Cast or heat chemical ranges and limits, %(a) UNS SAE-AISI number

number

C max

Mn

P

S

Pb

G12110

1211

0.13

0.60−0.90

0.07−012

0.10−0.15

...

G12120

1212

0.13

0.70−1.00

0.07−0.12

0.16−0.23

...

G12130

1213

0.13

0.70−1.00

0.07−0.12

0.24−0.33

...

G12150

1215

0.09

0.75−1.05

0.04−0.09

0.26−0.35

...

G12144 12L14 0.15 0.85−1.15 0.04−0.09 0.26−0.35 0.15−0.35 (a) When lead ranges or limits are required, the values in Table 1 apply. It is not common practice to produce the 12xx series of steels to specified limits for silicon because of its adverse effect on machinability. Source: Ref 1

Table 17 High-manganese carbon steel compositions Applicable only to semifinished products for forging, hot-rolled and cold-finished bars, wire rods, and seamless tubing Cast or heat chemical ranges and limits, %(a) UNS SAE-AISI number

number

C

Mn

P max

S max

G15130

1513

0.10−0.16

1.10−1.40

0.040

0.050

G15220

1522

0.18−0.24

1.10−1.40

0.040

0.050

G15240

1524

0.19−0.25

1.35−1.65

0.040

0.050

G15260

1526

0.22−0.29

1.10−1.40

0.040

0.050

G15270

1527

0.22−0.29

1.20−1.50

0.040

0.050

G15360

1536

0.30−0.37

1.20−1.50

0.040

0.050

G15410

1541

0.36−0.44

1.35−1.65

0.040

0.050

G15480

1548

0.44−0.52

1.10−1.40

0.040

0.050

G15510

1551

0.45−0.56

0.85−1.15

0.040

0.050


Page 252



G15520

1552

G15610

1561

01 Sep 2005

1.20−1.50

0.040

0.050

0.75−1.05

0.040

0.050

G15660 1566 0.040 0.60−0.71 0.85−1.15 (a) When silicon, lead, and boron ranges or limits are required, the values in Tables 1 and 2 apply. Source: Ref 1

0.050

0.47−0.55 0.55−0.65

Table 18 High-manganese carbon steel compositions Applicable only to structural shapes, plates, strip, sheets, and welded tubing Cast or heat chemical ranges and limits, %(a) UNS SAE-AISI number

number

C

Mn

P max

S max

Former SAE number

G15240

1524

0.18−0.25

1.30−1.65

0.040

0.050

1024

G15270

1527

0.22−0.29

1.20−1.55

0.040

0.050

1027

G15360

1536

0.30−0.38

1.20−1.55

0.040

0.050

1036

G15410

1541

0.36−0.45

1.30−1.65

0.040

0.050

1041

G15480

1548

0.43−0.52

1.05−1.40

0.040

0.050

1048

0.050

1052

G15520 1552 0.040 0.46−0.55 1.20−1.55 (a) When silicon ranges or limits are required, the values shown in Table 2 apply. Source: Ref 1

Table 19 Low-alloy steel compositions applicable to billets, blooms, slabs, and hot-rolled and cold-finished bars Slightly wider ranges of compositions apply to plates. The article "Carbon and Low-Alloy Steel Plate" in this volume lists SAE-AISI plate compositions Ladle chemical compositions limits, %(a) UNS SAE-AIS numb I Corresponding er number AISI number C Mn P S Si Ni Cr Mo V G133 00

1330

1330

0.28−0. 1.60−1. 0.035 33 90

0.040

0.15−0.35

...

...

...

...

G133 50

1335

1335

0.33−0. 1.60−1. 0.035 38 90

0.040

0.15−0.35

...

...

...

...

G134 00

1340

1340

0.38−0. 1.60−1. 0.035 43 90

0.040

0.15−0.35

...

...

...

...

G134 50

1345

1345

0.43−0. 1.60−1. 0.035 48 90

0.040

0.15−0.35

...

...

...

...

G402 30

4023

4023

0.20−0. 0.70−0. 0.035 25 90

0.040

0.15−0.35

...

...

G402 40

4024

4024

0.20−0. 0.70−0. 0.035 0.035−0.0 0.15−0.35 25 90 50

...

...

0.20−0.30

...

G402 70

4027

4027

0.25−0. 0.70−0. 0.035 30 90

0.15−0.35

...

...

0.20−0.30

...

G402 80

4028

4028

0.25−0. 0.70−0. 0.035 0.035−0.0 0.15−0.35 30 90 50

...

...

0.20−0.30

...

G403 20

4032

...

0.30−0. 0.70−0. 0.035 35 90

0.040

0.15−0.35

...

...

0.20−0.30

...

G403 70

4037

4037

0.35−0. 0.70−0. 0.035 40 90

0.040

0.15−0.35

...

...

0.20−0.30

...

G404 20

4042

...

0.40−0. 0.70−0. 0.035 45 90

0.040

0.15−0.35

...

...

0.20−0.30

...

G404 70

4047

4047

0.45−0. 0.70−0. 0.035 50 90

0.040

0.15−0.35

...

...

0.20−0.30

...

G411 80

4118

4118

0.18−0. 0.70−0. 0.035 23 90

0.040

0.15−0.35

...

0.40−0.60 0.08−015

...

G413 00

4130

4130

0.28−0. 0.40−0. 0.035 33 60

0.040

0.15−0.35

...

0.80−1.10 0.15−0.25

...

G413 50

4135

...

0.33−0. 0.70−0. 0.035 38 90

0.040

0.15−0.35

...

0.80−1.10 0.15−0.25

...

G413

4137

4137

0.35−0. 0.70−0. 0.035

0.040

0.15−0.35

...

0.80−1.10 0.15−0.25

...


0.040

Page 253



70

40

01 Sep 2005

90

G414 00

4140

4140

0.38−0. 0.75−1. 0.035 43 00

0.040

0.15−0.35

...

0.80−1.10 0.15−0.25

...

G414 20

4142

4142

0.40−0. 0.75−1. 0.035 45 00

0.040

0.15−0.35

...

0.80−1.10 0.15−0.25

...

G414 50

4145

4145

0.41−0. 0.75−1. 0.035 48 00

0.040

0.15−0.35

...

0.80−1.10 0.15−0.25

...

G414 70

4147

4147

0.45−0. 0.75−1. 0.035 50 00

0.040

0.15−0.35

...

0.80−1.10 0.15−0.25

...

G415 00

4150

4150

0.48−0. 0.75−1. 0.035 53 00

0.040

0.15−0.35

...

0.80−1.10 0.15−0.25

...

G416 10

4161

4161

0.56−0. 0.75−1. 0.035 64 00

0.040

0.15−0.35

...

0.70−0.90 0.25−0.35

...

G432 00

4320

4320

0.17−0. 0.45−0. 0.035 22 65

0.040

0.15−0.35 1.65−2.00 0.40−0.60 0.20−0.30

...

G434 00

4340

4340

0.38−0. 0.60−0. 0.035 43 80

0.040

0.15−0.35 1.65−2.00 0.70−0.90 0.20−0.30

...

E4340

0.38−0. 0.65−0. 0.025 43 85

0.025

0.15−0.35 1.65−2.00 0.70−0.90 0.20−0.30

...

G434 E4340(b) 06 G442 20

4422

...

0.20−0. 0.70−0. 0.035 25 90

0.040

0.15−0.35

...

...

0.35−0.45

...

G442 70

4427

...

0.24−0. 0.70−0. 0.035 29 90

0.040

0.15−0.35

...

...

0.35−0.45

...

G461 50

4615

4615

0.13−0. 0.45−0. 0.035 18 65

0.040

0.15−0.25 1.65−2.00

...

0.20−0.30

...

G461 70

4617

...

0.15−0. 0.45−0. 0.035 20 65

0.040

0.15−0.35 1.65−2.00

...

0.20−0.30

...

G462 00

4620

4620

0.17−0. 0.45−0. 0.035 22 65

0.040

0.15−0.35 1.65−2.00

...

0.20−0.30

...

G462 60

4626

4626

0.24−0. 0.45−0. 0.035 0.04 max 0.15−0.35 0.70−1.00 29 65

...

0.15−0.25

...

G471 80

4718

4718

0.16−0. 0.70−0. 21 90

0.90−1.20 0.35−0.55 0.30−0.40

...

G472 00

4720

4720

0.17−0. 0.50−0. 0.035 22 70

0.040

0.15−0.35 0.90−1.20 0.35−0.55 0.15−0.25

...

G481 50

4815

4815

0.13−0. 0.40−0. 0.035 18 60

0.040

0.15−0.35 3.25−3.75

...

0.20−0.30

...

G481 70

4817

4817

0.15−0. 0.40−0. 0.035 20 60

0.040

0.15−0.35 3.25−3.75

...

0.20−0.30

...

G482 00

4820

4820

0.18−0. 0.50−0. 0.035 23 70

0.040

0.15−0.35 3.25−3.75

...

0.20−0.30

...

G504 50B40(c) 01

...

0.38−0. 0.75−1. 0.035 43 00

0.040

0.15−0.35

...

0.40−0.60

...

...

G504 50B44(c) 41

50B44

0.43−0. 0.75−1. 0.035 48 00

0.040

0.15−0.35

...

0.40−0.60

...

...

...

0.43−0. 0.75−1. 0.035 48 00

0.040

0.15−0.35

...

0.20−0.35

...

...

G504 50B46(c) 61

50B46

0.44−0. 0.75−1. 0.035 49 00

0.040

0.15−0.35

...

0.20−0.35

...

...

G505 50B50(c) 01

50B50

0.48−0. 0.75−1. 0.035 53 00

0.040

0.15−0.35

...

0.40−0.60

...

...

...

0.56−0. 0.75−1. 0.035 64 00

0.040

0.15−0.35

...

0.40−0.60

...

...

50B60

0.56−0. 0.75−1. 0.035 64 00

0.040

0.15−0.35

...

0.40−0.60

...

...

G504 60

G506 00

5046

5060

G506 50B60(c) 01


...

...

...

Page 254



01 Sep 2005

G511 50

5115

...

0.13−0. 0.70−0. 0.035 18 90

0.040

0.15−0.35

...

0.70−0.90

...

...

G511 70

5117

5117

0.15−0. 0.70−0. 0.040 20 90

0.040

0.15−0.35

...

0.70−0.90

...

...

G512 00

5120

5120

0.17−0. 0.70−0. 0.035 22 90

0.040

0.15−0.35

...

0.70−0.90

...

...

G513 00

5130

5130

0.28−0. 0.70−0. 0.035 33 90

0.040

0.15−0.35

...

0.80−1.10

...

...

G513 20

5132

5132

0.30−0. 0.60−0. 0.035 35 80

0.040

0.15−0.35

...

0.75−1.00

...

...

G513 50

5135

5135

0.33−0. 0.60−0. 0.035 38 80

0.040

0.15−0.35

...

0.80−1.05

...

...

G514 00

5140

5140

0.38−0. 0.70−0. 0.035 43 90

0.040

0.15−0.35

...

0.70−0.90

...

...

G514 70

5147

5147

0.46−0. 0.70−0. 0.035 51 95

0.040

0.15−0.35

...

0.85−1.15

...

...

G515 00

5150

5150

0.48−0. 0.70−0. 0.035 53 90

0.040

0.15−0.35

...

0.70−0.90

...

...

G515 50

5155

5155

0.51−0. 0.70−0. 0.035 59 90

0.040

0.15−0.35

...

0.70−0.90

...

...

G516 00

5160

5160

0.56−0. 0.75−1. 0.035 64 00

0.040

0.15−0.35

...

0.70−0.90

...

...

51B60

0.56−0. 0.75−1. 0.035 64 00

0.040

0.15−0.35

...

0.70−0.90

...

...

G516 51B60(c) 01 G509 86

50100(b)

...

0.98−1. 0.25−0. 0.025 10 45

0.025

0.15−0.35

...

0.40−0.60

...

...

G519 86

51100(b)

E51100

0.98−1. 0.25−0. 0.025 10 45

0.025

0.15−0.35

...

0.90−1.15

...

...

G529 86

52100(b)

E52100

0.98−1. 0.25−0. 0.025 10 45

0.025

0.15−0.35

...

1.30−1.60

...

...

G611 80

6118

6118

0.16−0. 0.50−0. 0.035 21 70

0.040

0.15−0.35

...

0.50−0.70

...

0.10−0.15

G615 00

6150

6150

0.48−0. 0.70−0. 0.035 53 90

0.040

0.15−0.35

...

0.80−1.10

...

0.15 min

G811 50

8115

8115

0.13−0. 0.70−0. 0.035 18 90

0.040

0.15−0.35 0.20−0.40 0.30−0.50 0.08−0.15

...

81B45

0.43−0. 0.75−1. 0.035 48 00

0.040

0.15−0.35 0.20−0.40 0.35−0.55 0.08−0.15

...

G814 81B45(c) 51 G861 50

8615

8615

0.13−0. 0.70−0. 0.035 18 90

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G861 70

8617

8617

0.15−0. 0.70−0. 0.035 20 90

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G862 00

8620

8620

0.18−0. 0.70−0. 0.035 23 90

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G862 20

8622

8622

0.20−0. 0.70−0. 0.035 25 90

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G862 50

8625

8625

0.23−0. 0.70−0. 0.035 28 90

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G862 70

8627

8627

0.25−0. 0.70−0. 0.035 30 90

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G863 00

8630

8630

0.28−0. 0.70−0. 0.035 33 90

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G863 70

8637

8637

0.35−0. 0.75−1. 0.035 40 00

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G864

8640

8640

0.38−0. 0.75−1. 0.035

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...


Page 255



00

43

01 Sep 2005

00

G864 20

8642

8642

0.40−0. 0.75−1. 0.035 45 00

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G864 50

8645

8645

0.43−0. 0.75−1. 0.035 48 00

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G864 86B45(c) 51

...

0.43−0. 0.75−1. 0.035 48 00

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G865 00

8650

...

0.48−0. 0.75−1. 0.035 53 00

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G865 50

8655

8655

0.51−0. 0.75−1. 0.035 59 00

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G866 00

8600

...

0.56−0. 0.75−1. 0.035 64 00

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.15−0.25

...

G872 00

8720

8720

0.18−0. 0.70−0. 0.035 23 90

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.20−0.30

...

G874 00

8740

8740

0.38−0. 0.75−1. 0.035 43 00

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.20−0.30

...

G882 20

8822

8822

0.20−0. 0.75−1. 0.035 25 00

0.040

0.15−0.35 0.40−0.70 0.40−0.60 0.30−0.40

...

G925 40

9254

...

0.51−0. 0.60−0. 0.035 59 80

0.040

1.20−1.60

...

0.60−0.80

...

...

G926 00

9260

9260

0.56−0. 0.75−1. 0.035 64 00

0.040

1.80−2.20

...

...

...

...

G931 06

9310(b)

...

0.08−0. 0.45−0. 0.025 13 65

0.025

0.15−0.35 3.00−3.50 1.00−1.40 0.08−0.15

...

G941 94B15(c) 51

...

0.13−0. 0.75−1. 0.035 18 00

0.040

0.15−0.35 0.30−0.60 0.30−0.50 0.08−0.15

...

G941 94B17(c) 71

94B17

0.15−0. 0.75−1. 0.035 20 00

0.040

0.15−0.35 0.30−0.60 0.30−0.50 0.08−0.15

...

G943 94B30(c) 94B30 0.040 0.15−0.35 0.30−0.60 0.30−0.50 0.08−0.15 ... 0.28−0. 0.75−1. 0.035 01 33 00 (a) Small quantities of certain elements that are not specified or required may be found in alloy steels. These elements are to be considered as incidental and are acceptable to the following maximum amount: copper to 0.35%, nickel to 0.25%, chromium to 0.20%, and molybdenum to 0.06%. (b) Electric furnace steel. (c) Boron content is 0.0005−0.003%. Source: Ref 16

Table 20 SAE potential standard steel compositions SAE PS number(a) PS 10

Ladle chemical composition limits, wt% C

Mn

P max

S max

Si

Ni

Cr

Mo

B

0.19−0.24

0.95−1.25

0.035

0.040

0.15−0.35

0.20−0.40

0.25−0.40

0.05−0.10

...

0.040

0.15−0.35

...

0.40−0.60

0.13−0.20

...

PS 15

0.18−0.23

0.90−1.20

0.035

PS 16

0.20−0.25

0.90−1.20

0.035

0.040

0.15−0.35

...

0.40−0.60

0.13−0.20

...

PS 17

0.23−0.28

0.90−1.20

0.035

0.040

0.15−0.35

...

0.40−0.60

0.13−0.20

...

PS 18

0.25−0.30

0.90−1.20

0.035

0.040

0.15−0.35

...

0.40−0.60

0.13−0.20

...

PS 19

0.18−0.23

0.90−1.20

0.035

0.040

0.15−0.35

...

0.40−0.60

0.08−0.15

0.0005−0.003

0.040

0.15−0.35

...

0.40−0.60

0.13−0.20

...

PS 20

0.13−0.18

0.90−1.20

0.035

PS 21

0.15−0.20

0.90−1.20

0.035

0.040

0.15−0.35

...

0.40−0.60

0.13−0.20

...

PS 24

0.18−0.23

0.75−1.00

0.035

0.040

0.15−0.35

...

0.45−0.65

0.20−0.30

...

PS 30

0.13−0.18

0.70−0.90

0.035

0.040

0.15−0.35

0.70−1.00

0.45−0.65

0.45−0.60

...

PS 31

0.15−0.20

0.70−0.90

0.035

0.040

0.15−0.35

0.70−1.00

0.45−0.65

0.45−0.60

...

0.040

0.15−0.35

0.70−1.00

0.45−0.65

0.45−0.60

...

PS 32

0.18−0.23

0.70−0.90

0.035

PS 33(b)

0.17−0.24

0.85−1.25

0.035

0.040

0.15−0.35

0.20 min

0.20 min

0.05 min

...

PS 34

0.28−0.33

0.90−1.20

0.035

0.040

0.15−0.35

...

0.40−0.60

0.13−0.20

...


Page 256



01 Sep 2005

PS 36

0.38−0.43

0.90−1.20

0.035

0.040

0.15−0.35

...

0.45−0.65

0.13−0.20

...

PS 38

0.43−0.48

0.90−1.20

0.035

0.040

0.15−0.35

...

0.45−0.65

0.13−0.20

...

PS 39

0.48−0.53

0.90−1.20

0.035

0.040

0.15−0.35

...

0.45−0.65

0.13−0.20

...

PS 40

0.51−0.59

0.90−1.20

0.035

0.040

0.15−0.35

...

0.45−0.65

0.13−0.20

...

PS 54

0.19−0.25

0.70−1.05

0.035

0.040

0.15−0.35

...

0.40−0.70

0.05 min

...

0.040

PS 55

0.15−0.20

0.70−1.00

0.035

0.15−0.35

1.65−2.00

0.45−0.65

0.65−0.80

...

PS 56

0.080−0.13

0.70−1.00

0.035

0.040

0.15−0.35

1.65−2.00

0.445−0.65

0.65−0.80

...

PS 57

0.08 max

1.25 max

0.040

0.15−0.35

1.00 max

...

17.00−19.00

1.75−2.25

...

PS 58

0.16−0.21

1.00−1.30

0.035

0.040

0.15−0.35

...

0.45−0.65

...

...

0.040

0.15−0.35

...

0.70−0.90

...

...

PS 59

0.18−0.23

1.00−1.30

0.035

PS 61

0.23−0.28

1.00−1.30

0.035

0.040

0.15−0.35

...

0.70−0.90

...

...

PS 63

0.31−0.38

0.75−1.10

0.035

0.040

0.15−0.35

...

0.45−0.65

...

0.0005−0.003

PS 64

0.16−0.21

1.00−1.30

0.035

0.040

0.15−0.35

...

0.70−0.90

...

...

PS 65

0.21−0.26

1.00−1.30

0.035

0.040

0.15−0.35

...

0.70−0.90

...

...

0.40−0.70

0.035

0.040

0.08−0.15

...

PS 66(c)

0.16−0.21

0.15−0.35

1.65−2.00

0.45−0.75

PS 67 0.035 0.040 ... ... 0.42−0.49 0.80−1.20 0.15−0.35 0.85−1.20 0.25−0.35 (a) Some PS steels may be supplied to a hardenability requirement. (b) Supplied to a hardenability requirement of 15 HRC points within the range of 23−43 HRC at J4 (4=16 in. distance from quenched end), subject to agreement between producer and user. (c) PS 66 a vanadium content of 0.10−0.15%. Source: Ref 17

Table 21 Composition ranges and limits for SAE HSLA steels Heat composition limits, %(a)

SAE designation(b)

C max

Mn max

P max

942X

0.21

1.35

0.04

945A

0.15

1.00

0.04

945C

0.23

1.40

0.04

945X

0.22

1.35

0.04

950A

0.15

1.30

0.04

950B

0.22

1.30

0.04

950C

0.25

1.60

0.04

950D

0.15

1.00

0.15

950X

0.23

1.35

0.04

955X

0.25

1.35

0.04

960X

0.26

1.45

0.04

965X

0.26

1.45

0.04

970X

0.26

1.65

0.04

980X 0.26 1.65 0.04 (a) Maximum contents of sulfur and silicon for all grades: 0.050% S, 0.90% Si. (b) Second and third digits of designation indicate minimum yield strength in ksi. Suffix X indicates that the steel contains niobium, vanadium, nitrogen, or other alloying elements. A second suffix K indicates that the steel is produced fully killed using fine-grain practice; otherwise, the steel is produced semikilled. Source: Ref 18

Table 22 Composition ranges and limits for former standard SAE steels Composition, wt%

SAE numb er

AISI number

UNS number

1009

1009

...

1011

...

1033

1033

C 0.15 max

Mn

P max(b) S max(b)

0.60 max

Si

Cr

Ni

Mo

Date of obsolesc V min ence

0.040

0.050

...

...

...

...

...

1965

G10110

0.08−0. 0.60−0. 13 90

0.040

0.050

...

...

...

...

...

1977

...

0.30−0. 0.70−1. 36 00

0.040

0.050

...

...

...

...

...

1965


Page 257



01 Sep 2005

1034

C1034

...

0.32−0. 0.50−0. 38 80

0.040

0.050

...

...

...

...

...

1968

1059( a)

...

...

0.55−0. 0.50−0. 65 80

0.040

0.050

...

...

...

...

...

1968

1062

C1062

...

0.54−0. 0.85−1. 65 15

0.040

0.050

...

...

...

...

...

1953

1086( a)

...

G10860

0.80−0. 0.30−0. 94 50

0.040

0.050

...

...

...

...

...

1977

1109

1109

G11090

0.08−0. 0.60−0. 13 90

0.040

0.08−0.13

...

...

...

...

...

1977

1111

B1111

...

0.13 max

0.60−0. 0.07−0.12 0.10−0.15 90

...

...

...

...

...

1969

1112

B1112

...

0.13 max

0.70−1. 0.07−0.12 0.16−0.23 00

...

...

...

...

...

1969

1113

B1113

...

0.13 max

0.70−1. 0.07−0.12 0.24−0.33 00

...

...

...

...

...

1969

1114

C1114

...

0.10−0. 1.00−1. 16 30

0.040

0.08−0.13

...

...

...

...

...

1952

1115

1115

...

0.13−0. 0.60−0. 18 90

0.040

0.08−0.13

...

...

...

...

...

1965

1116

C1116

...

0.14−0. 1.10−1. 20 40

0.040

0.16−0.23

...

...

...

...

...

1952

1119

1119

G11190

0.14−0. 1.00−1. 20 30

0.040

0.24−0.23

...

...

...

...

...

1977

1120

1120

...

0.18−0. 0.70−1. 23 00

0.040

0.08−0.13

...

...

...

...

...

1965

1126

1126

...

0.23−0. 0.70−1. 29 00

0.040

0.08−0.13

...

...

...

...

...

1965

1132

1132

G11320

0.27−0. 1.35−1. 34 65

0.040

0.08−0.13

...

...

...

...

...

1977

1138

1138

...

0.34−0. 0.70−1. 40 00

0.040

0.08−0.13

...

...

...

...

...

1965

1145

1145

G11450

0.42−0. 0.70−1. 49 00

0.040

0.04−0.07

...

...

...

...

...

1977

1320

A1320

...

0.18−0. 1.60−1. 23 90

0.040

0.040

0.20−0.35

...

...

...

...

1956

1518

...

G15180

0.15−0. 1.10−1. 21 40

0.040

0.050

...

...

...

...

...

1977

1525

...

G15250

0.23−0. 0.80−1. 29 10

0.040

0.050

...

...

...

...

...

1977

1547

...

G15470

0.43−0. 1.35−1. 51 65

0.040

0.050

...

...

...

...

...

1977

1572

...

G15720

0.65−0. 1.00−1. 76 30

0.040

0.050

...

...

...

...

...

1977

2317

A2317

...

0.15−0. 0.40−0. 20 60

0.040

0.040

0.20−0.35

...

3.25−3.75

...

...

1956

2330

A2330

...

0.28−0. 0.60−0. 33 80

0.040

0.040

0.20−0.35

...

3.25−3.75

...

...

1953

2340

A2340

...

0.38−0. 0.70−0. 43 90

0.040

0.040

0.20−0.35

...

3.25−3.75

...

...

1953

2345

A2345

...

0.43−0. 0.70−0. 48 90

0.040

0.040

0.20−0.35

...

3.25−3.75

...

...

1952

2512

E2512

...

0.09−0. 0.45−0. 14 60

0.025

0.025

0.20−0.35

...

4.75−5.25

...

...

1953

2515

A2515

...

0.12−0. 0.40−0.

0.040

0.040

0.20−0.35

...

4.75−5.25

...

...

1956


Page 258



17

01 Sep 2005

60

2517

E2517

...

0.15−0. 0.45−0. 20 60

0.025

0.025

0.20−0.35

4.75−5.25

...

...

1959

3115

A3115

...

0.13−0. 0.40−0. 18 60

0.040

0.040

0.20−0.35 0.55−0.75 1.10−1.40

...

...

1953

3120

A3120

...

0.17−0. 0.60−0. 22 80

0.040

0.040

0.20−0.35 0.55−0.75 1.10−1.40

...

...

1956

3130

A3130

...

0.28−0. 0.60−0. 33 80

0.040

0.040

0.20−0.35 0.55−0.75 1.10−1.40

...

...

1956

3135

3135

...

0.33−0. 0.60−0. 38 80

0.040

0.040

0.20−0.35 0.55−0.75 1.10−1.40

...

...

1960

X314 0

A3141

...

0.38−0. 0.70−0. 43 90

0.040

0.040

0.20−0.35 0.70−0.90 1.10−1.40

...

...

1947

3140

3140

...

0.38−0. 0.70−0. 43 90

0.040

0.040

0.20−0.35 0.55−0.75 1.10−1.40

...

...

1964

3145

A3145

...

0.43−0. 0.70−0. 48 90

0.040

0.040

0.20−0.35 0.70−0.90 1.10−1.40

...

...

1952

3150

A3150

...

0.48−0. 0.70−0. 53 90

0.040

0.040

0.20−0.35 0.70−0.90 1.10−1.40

...

...

1952

3215

...

...

0.10−0. 0.30−0. 20 60

0.040

0.050

0.15−0.30 0.90−1.25 1.50−2.00

...

...

1941

3220

...

...

0.15−0. 0.30−0. 25 60

0.040

0.050

0.15−0.30 0.90−1.25 1.50−2.00

...

...

1941

3230

...

...

0.25−0. 0.30−0. 35 60

0.040

0.050

0.15−0.30 0.90−1.25 1.50−2.00

...

...

1941

3240

A3240

...

0.35−0. 0.30−0. 45 60

0.040

0.040

0.15−0.30 0.90−1.25 1.50−2.00

...

...

1941

3245

...

...

0.40−0. 0.30−0. 50 60

0.040

0.040

0.15−0.30 0.90−1.25 1.50−2.00

...

...

1941

3250

...

...

0.45−0. 0.30−0. 55 60

0.040

0.040

0.15−0.30 0.90−1.25 1.50−2.00

...

...

1941

3310

E3310

...

0.08−0. 0.45−0. 13 60

0.025

0.025

0.20−0.35 0.40−1.75 3.25−3.75

...

...

1964

3312

...

...

0.08−0. 0.40−0. 13 60

0.025

0.025

0.20−0.35 0.140−1.7 3.25−3.75 5

...

...

1948

3316

E3316

...

0.14−0. 0.45−0. 19 60

0.025

0.025

0.20−0.35 1.40−1.75 3.25−3.75

...

...

1956

3325

...

...

20−30

0.30−0. 60

0.040

0.050

0.15−0.30 1.25−1.75 3.25−3.75

...

...

1936

3335

...

...

30−40

0.30−0. 60

0.040

0.050

0.15−0.30 1.25−1.75 3.25−3.75

...

...

1936

3340

...

...

35−45

0.30−0. 60

0.040

0.050

0.15−0.30 1.25−1.75 3.25−3.75

...

...

1936

3415

...

...

0.10−0. 0.30−0. 20 60

0.040

0.050

0.15−0.30 0.60−0.95 2.75−3.25

...

...

1941

3435

...

...

0.30−0. 0.30−0. 40 60

0.040

0.050

0.15−0.30 0.60−0.95 2.75−3.25

...

...

1936

3450

...

...

0.45−0. 0.30−0. 55 60

0.040

0.050

0.15−0.30 0.60−0.95 2.75−3.25

...

...

1936

4012

4012

G40120

0.09−0. 0.75−1. 14 00

0.035

0.040

0.15−0.30

...

...

0.15−0.25

...

1977

4053

4053

...

0.50−0. 0.75−1. 56 00

0.040

0.040

0.20−0.35

...

...

0.20−0.30

...

1956

4063

4063

G40630

0.60−0. 0.75−1. 67 00

0.040

0.040

0.20−0.35

...

...

0.20−0.30

...

1964


...

Page 259



4068

A4068

...

0.63−0. 0.75−1. 70 00

0.040

0.040

0.20−0.35

4119

A4119

...

0.17−0. 0.70−0. 22 90

0.040

0.040

4125

A4125

...

0.23−0. 0.70−0. 28 90

0.040

4317

4317

...

0.15−0. 0.45−0. 20 65

4337

4337

G43370

4419

4520

4419 H 4608

...

0.20−0.30

...

1957

0.20−0.35 0.40−0.60

...

0.20−0.30

...

1956

0.040

0.20−0.35 0.40−0.60

...

0.20−0.30

...

1950

0.040

0.040

0.20−0.35 0.40−0.60 1.65−2.00 0.20−0.30

...

1953

0.35−0. 0.60−0. 40 80

0.040

0.040

0.20−0.35 0.70−0.90 1.65−2.00 0.20−0.30

...

1964

...

0.18−0. 0.45−0. 23 65

0.035

0.040

0.15−0.30

...

...

0.45−0.60

...

1977

4419H

...

0.17−0. 0.35−0. 23 75

0.035

0.040

0.15−0.30

...

...

0.45−0.60

...

1977

4608

...

0.06−0. 0.25−0. 11 45

0.040

0.040

0.25 max

...

1.40−1.75 0.15−0.25

...

1956

0.10−0. 0.45−0. 15 65

0.040

0.040

0.20−0.35

...

1.65−2.00 0.20−0.30

...

1957

...

0.18−0. 0.50−0. 23 70

0.040

0.040

0.20−0.35

...

1.65−2.00 0.20−0.30

...

1956

G46210

0.18−0. 0.70−0. 23 90

0.035

0.040

0.15−0.30

...

1.65−2.00 0.20−0.30

...

1977

46B12 46B12(c) (c)

...

01 Sep 2005

X462 0

X4620

4621

4621

4621 H

4621H

...

0.17−0. 0.60−1. 23 00

0.035

0.040

0.15−0.30

...

1.55−2.00 0.20−0.30

...

1977

4640

A4640

...

0.38−0. 0.60−0. 43 80

0.040

0.040

0.20−0.35

...

1.65−2.00 0.20−0.30

...

1952

4812

4817

...

0.10−0. 0.40−0. 15 60

0.040

0.040

0.20−0.35

...

3.25−3.75 0.20−0.30

...

1956

5015

5015

G50150

0.12−0. 0.030−0 17 .50

0.035

0.040

0.15−0.30 0.30−0.50

...

...

...

1977

5045

5045

...

0.43−0. 0.70−0. 48 90

0.040

0.040

0.20−0.35 0.55−0.75

...

...

...

1953

5145

5145

G51450

0.43−0. 0.70−0. 48 90

0.035

0.040

0.15−0.30 0.70−0.90

...

...

...

1977

5145 H

5145H

H51450

0.42−0. 0.60−1. 49 00

0.035

0.040

0.15−0.30 0.60−1.00

...

...

...

1977

5152

5152

...

0.48−0. 0.70−0. 55 90

0.040

0.040

0.20−0.35 0.90−1.20

...

...

...

1956

6115

...

...

0.10−0. 0.30−0. 20 60

0.040

0.050

0.15−0.30 0.80−0.10

...

...

0.15

1936

6117

6117

...

0.15−0. 0.70−0. 20 90

0.040

0.040

0.20−0.35 0.70−0.90

...

...

0.10

1956

6120

6120

...

0.17−0. 0.70−0. 22 90

0.040

0.040

0.20−0.35 0.70−0.90

...

...

0.10

1961

6125

...

...

0.20−0. 0.60−0. 30 90

0.040

0.050

0.15−0.30 0.80−0.10

...

...

0.15

1936

6130

...

...

0.25−0. 0.60−0. 35 90

0.040

0.050

0.15−0.30 0.80−0.10

...

...

0.15

1936

6135

...

...

0.30−0. 0.60−0. 40 90

0.040

0.050

0.15−0.30 0.80−0.10

...

...

0.15

1941

6140

...

...

0.35−0. 0.60−0. 45 90

0.040

0.050

0.15−0.30 0.80−0.10

...

...

0.15

1936

6145

6145

...

0.43−0. 0.70−0. 48 90

0.040

0.050

0.20−0.35 0.80−0.10

...

...

0.15

1956

6195

...

...

0.90−1. 0.20−0.

0.030

0.035

0.15−0.30 0.80−0.10

...

...

0.15

1936


Page 260



05

01 Sep 2005

45

71360 (d)

...

...

0.50−0. 0.30 max 70

0.035

0.040

0.15−0.30 3.00−4.00

...

...

...

1936

71660 (e)

...

...

0.50−0. 0.30 max 70

0.035

0.040

0.15−0.30 3.00−4.00

...

...

...

1936

7260(f )

...

...

0.50−0. 0.30 max 70

0.035

0.040

0.15−0.30 0.50−1.00

...

...

...

1936

8632

8632

...

0.30−0. 0.70−0. 35 90

0.040

0.040

0.20−0.35 0.40−0.60 0.40−0.70 0.15−0.25

...

1951

8635

8635

...

0.33−0. 0.75−1. 38 00

0.040

0.040

0.20−0.35 0.40−0.60 0.40−0.70 0.15−0.25

...

1956

8641

8641

...

0.38−0. 0.75−1. 43 00

0.040

0.040−0.0 0.20−0.35 0.40−0.60 0.40−0.70 0.15−0.25 60

...

1956

8653

8653

...

0.50−0. 0.75−1. 56 00

0.040

0.040

0.20−0.35 0.50−0.80 0.40−0.70 0.15−0.25

...

1956

8647

8647

...

0.45−0. 0.75−1. 50 00

0.040

0.040

0.20−0.35 0.40−0.60 0.40−0.70 0.15−0.25

...

1948

8715

8715

...

0.13−0. 0.70−0. 18 90

0.040

0.040

0.20−0.35 0.40−0.60 0.40−0.70 0.20−0.30

...

1956

8717

8717

...

0.15−0. 0.70−0. 20 90

0.040

0.040

0.20−0.35 0.40−0.60 0.40−0.70 0.20−0.30

...

1956

8719

8719

...

0.18−0. 0.60−0. 23 80

0.040

0.040

0.20−0.35 0.40−0.60 0.40−0.70 0.20−0.30

...

1952

8735

8735

G87350

0.33−0. 0.75−1. 38 00

0.040

0.040

0.20−0.35 0.40−0.60 0.40−0.70 0.20−0.30

...

1952

8742

8742

G87420

0.40−0. 0.75−1. 45 00

0.040

0.040

0.20−0.35 0.40−0.60 0.40−0.70 0.20−0.30

...

1964

8745

8745

...

0.43−0. 0.75−1. 48 00

0.040

0.040

0.20−0.35 0.40−0.60 0.40−0.70 0.20−0.30

...

1953

8750

8750

...

0.48−0. 0.75−1. 53 00

0.040

0.040

0.20−0.35 0.40−0.60 0.40−0.70 0.20−0.30

...

1956

9250

9250

...

0.45−0. 0.60−0. 55 90

0.040

0.040

1.80−2.20

...

...

...

...

1941

9255

9255

G92550

0.51−0. 0.70−0. 59 95

0.035

0.040

1.80−2.20

...

...

...

...

1977

9261

9261

...

0.55−0. 0.75−1. 65 00

0.040

0.040

1.80−2.20 0.10−0.25

...

...

...

1956

9262

9262

G92620

0.55−0. 0.75−1. 65 00

0.040

0.040

1.80−2.20 0.25−0.40

...

...

...

1961

9315

E9315

...

0.13−0. 0.45−0. 18 65

0.025

0.025

0.20−0.35 1.00−1.40 3.00−3.50 0.08−0.15

...

1959

9317

E9317

...

0.15−0. 0.45−0. 20 65

0.025

0.025

0.20−0.35 1.00−1.40 3.00−3.50 0.08−0.15

...

1959

9437

9437

...

0.35−0. 0.90−1. 40 20

0.040

0.040

0.20−0.35 0.30−0.50 0.30−0.60 0.08−0.15

...

1950

9440

9440

...

0.38−0. 0.90−1. 43 20

0.040

0.040

0.20−0.35 0.30−0.50 0.30−0.60 0.08−0.15

...

1950

G94401

0.38−0. 0.75−1. 43 00

0.040

0.040

0.20−0.35 0.30−0.50 0.30−0.60 0.08−0.15

...

1964

94B40 (c)

94B40

9442

9442

...

0.40−0. 0.90−1. 45 20

0.040

0.040

0.20−0.35 0.30−0.50 0.30−0.60 0.08−0.15

...

1950

9445

9445

...

0.43−0. 0.90−1. 48 20

0.040

0.040

0.20−0.35 0.30−0.50 0.30−0.60 0.08−0.15

...

1950

9447

9447

...

0.45−0. 0.90−1. 50 20

0.040

0.040

0.20−0.35 0.30−0.50 0.30−0.60 0.08−0.15

...

1950


Page 261



01 Sep 2005

9747

9747

...

0.45−0. 0.50−0. 50 80

0.040

0.040

0.20−0.35 0:10−0.25 0.40−0.70 0.15−0.25

...

1950

9763

9763

...

0.60−0. 0.50−0. 67 80

0.040

0.040

0.20−0.35 0.10−0.25 0.40−0.70 0.15−0.25

...

1950

9840

9840

G98400

0.38−0. 0.70−0. 43 90

0.040

0.040

0.20−0.35 0.70−0.90 0.85−1.15 0.20−0.30

...

1964

9845

9845

...

0.43−0. 0.70−0. 48 90

0.040

0.040

0.20−0.35 0.70−0.90 0.85−1.15 0.20−0.30

...

1950

9850

9850

G98500

0.48−0. 0.70−0. 53 90

0.040

0.040

0.20−0.35 0.70−0.90 0.85−1.15 0.20−0.30

...

1961

43BV 12(c)

...

...

0.08−0. 0.75−1. 13 00

...

...

0.20−0.35 0.40−0.60 1.65−2.00 0.20−0.30 0.03

...

43BV ... ... ... ... ... 0.10−0. 0.45−0. 0.20−0.35 0.40−0.60 1.65−2.00 0.08−0.15 0.03 14(c) 15 65 (a) These grades remain standard for wire rods. (b) Limits apply to semifinished products for forgings, bars, wire rods, and seamless tubing. (c) Boron content 0.0005−0.003%. (d) Contains 12.00−15.00% W. (e) Contains 15.00−18.00% W. (f) Contains 1.50−2.00% W. Source: Ref 19

Low-carbon steels contain up to 0.30% C. The largest category of this class of steel is flat-rolled products (sheet or strip) usually in the cold-rolled and annealed condition. The carbon content for these high-formability steels is very low, less than 0.10% C, with up to 0.4% Mn. Typical uses are in automobile body panels, tin plate, and wire products. For rolled steel structural plates and sections, the carbon content may be increased to approximately 0.30%, with higher manganese up to 1.5%. These latter materials may be used for stampings, forgings, seamless tubes, and boiler plate. Medium-carbon steels are similar to low-carbon steels except that the carbon ranges from 0.30 to 0.60% and the manganese from 0.60 to 1.65%. Increasing the carbon content to approximately 0.5% with an accompanying increase in manganese allows medium-carbon steels to be used in the quenched and tempered condition. The uses of medium carbon-manganese steels include shafts, couplings, crankshafts, axles, gears, and forgings. Steels in the 0.40 to 0.60% C range are also used for rails, railway wheels, and rail axles. High-carbon steels contain from 0.60 to 1.00% C with manganese contents ranging from 0.30 to 0.90%. High-carbon steels are used for spring materials and high-strength wires. Ultrahigh-carbon steels are experimental alloys containing approximately 1.25 to 2.0% C. These steels are thermomechanically processed to produce microstructures that consist of ultrafine, equiaxed grains of ferrite and a uniform distribution of fine, spherical, discontinuous proeutectoid carbide particles (Ref 13). Such microstructures in these steels have led to superplastic behavior (Ref 14). Properties of these experimental steels are described in Forming and Forging, Volume 14 of ASM Handbook, formerly 9th Edition Metals Handbook (see the Appendix to the article "Superplastic Sheet Forming," entitled "Superplasticity in Iron-Base Alloys").

High-Strength Low-Alloy Steels High-strength low-alloy (HSLA) steels, or microalloyed steels, are designed to provide better mechanical properties and/or greater resistance to atmospheric corrosion than conventional carbon steels. They are not considered to be alloy steels in the normal sense because they are designed to meet specific mechanical properties rather than a chemical composition (HSLA steels have yield strengths of more than 275 MPa, or 40 ksi). The chemical composition of a specific HSLA steel may vary for different product thickness to meet mechanical property requirements. The HSLA steels have low carbon contents (0.50 to ~0.25% C) in order to produce adequate formability and weldability, and they have manganese contents up to 2.0%. Small quantities of chromium, nickel, molybdenum, copper, nitrogen, vanadium, niobium, titanium, and zirconium are used in various combinations. The HSLA steels are commonly furnished in the as-rolled condition. They may also be supplied in a controlled-rolled, normalized, or precipitation-hardened condition to meet specific property requirements. Primary applications for HSLA steels include oil and gas line pipe, ships, offshore structures, automobiles, off-highway equipment, and pressure vessels. HSLA Classification. The types of HSLA steels commonly used include (Ref 15): • Weathering steels, designed to exhibit superior atmospheric corrosion resistance • Control-rolled steels, hot rolled according to a predetermined rolling schedule designed to develop a highly deformed austenite structure that will transform to a very fine equiaxed ferrite structure on cooling • Pearlite-reduced steels, strengthened by very fine-grain ferrite and precipitation hardening but with low carbon content and therefore little or no pearlite in the microstructure • Microalloyed steels, with very small additions (generally 300−1200 incl >12−48 incl

Coils and cut lengths

1.2−4.5

0.045−0.180 incl

>1200


A 569, A 621, or A 622

A 569M, A 621M, or A 622M

>48

A 569, A 621, or A 622

A 569M, A 621M, or A 622M

6.0−12.5 0.230−0.500 incl

>300−1200 incl >12−48 incl

Coils only

A 635

A 635M

4.5−12.5 0.180−0.500 incl

>1200−1800 incl

>48−72 incl

Coils only

A 635

A 635M

1.2−5.0

0.045−0.203 incl

≤200

≤6


A 569, A 621, or A 622

A 569M, A 621M, or A 622M

1.2−6.0

0.045−0.229 incl

>200−300 incl

>6−12 incl


A 569, A 621, or A 622

A 569M, A 621M, or A 622M

6.0−12.5 0.230−0.500 incl

>200−300 incl

>8−12 incl

Coils only

A 635

A 635M

>50−300 incl

>2−12 incl

(a)

A 366, A 619, or A 620

A 366M, A 619M, or A 620M

>0.014

>300

>12

(b)

A 366, A 619, or A 620

A 366M, A 619M, or A 620M

≤0.250

>12−600 incl

Cold-rolled 0.35−2.0 0.014−0.082 sheet incl ≥0.35 Cold-rolled ≤6.0 strip

(c) A 109 A 109M >0.50−23.9 incl (a) Incl, inclusive. (b) Cold-rolled sheet, coils, and cut lengths, slit from wider coils with cut edge (only), thicknesses 0.356−2.08 mm (0.014−0.082 in.) and 0.25% C (max) by cost analysis. (c) When no special edge or finish (other than matte, commercial bright, or luster) is required and/or single-strand rolling of widths under 610 mm (24 in.) is not required. (d) Width 51−305 mm (2−12 in.) with thicknesses of 0.356−2.08 mm (0.014−0.082 in.) are classified as sheet when slit from wider coils, have a cut edge only, and contain 0.25% C (max) by cost analysis. Source: Ref 2

Table 4(a) Summary of available types of hot-rolled and cold-rolled plain carbon steel sheet and strip Surface finish Temper-rolled;


Annealed last;

Page 337


Quality or temper

Carbon and Low-Alloy Steel Sheet and Str...

Applicable AISI-SAE grade basic specification designatio n number

for exposed parts(a) Description

Symbol

01 Sep 2005

Edge(b)

for unexposed parts(a) Description

Symbol

Descriptio n

Symbol

Hot-rolled sheet Commercial Quality

Drawing quality

A 569, A 635

A 621

Drawing quality, special killed A 622

1008−1012 As-rolled (black)

A

As-rolled (black)

A

Mill

M

Pickled⎯dry

P

Pickled⎯dry

P

Mill

M

Pickled and oiled

O

Pickled and oiled

O

Cut

C

A

As-rolled (black)

A

Mill

M

Pickled⎯dry

P

Pickled⎯dry

P

Mill

M

Pickled and oiled

O

Pickled and oiled

O

Cut

C

A

As-rolled (black)

A

Mill

M

Pickled⎯dry

P

Pickled⎯dry

P

Mill

M

Pickled and oiled

O

Pickled and oiled

O

Cut

C

A

As-rolled (black)

A

Mill

M

Pickled⎯dry

P

Pickled⎯dry

P

Mill

M

Pickled and oiled

O

Pickled and oiled

O

Cut

M

A

As-rolled (black)

A

Square

S

Pickled⎯dry

P

Pickled⎯dry

P

Square

S

Pickled and oiled

O

Pickled and oiled

O

Square

S

A

As-rolled (black)

A

Cut

C

Pickled⎯dry

P

Pickled⎯dry

P

Cut

C

Pickled and oiled

O

Pickled and oiled

O

Cut

C

E

Matte

U



Hot-rolled strip Commercial quality

Drawing quality

A 569

A 621

Drawing quality, special killed A 622




Cold-rolled sheet Commercial quality

A 366

1008−1012 Matte Commercial bright Luster

Drawing quality

A 619

1006−1008 Matte Commercial bright

(c)

Cut

(c)

L

Cut

(c)

Cut

(c)

B

Cut

(c)

L

Cut

(c)

Cut

(c)

B

Cut

(c)

Luster

L

Cut

(c)

Matte

1

Matte

1

(b)

1, 2, 3, 4, 5, 6

Regular bright

2

Regular bright

2

(b)

1, 2, 3, 4, 5, 6

Luster Drawing quality, special killed A 620

Cut

B

1006−1008 Matte Commercial bright

E

E

Matte

Matte

U

U

Cold-rolled strip Temper description numbers Ã1, 2, 3, 4, 5

A 109

(d)


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3

01 Sep 2005

Best bright

3

(b)

1, 2, 3, 4, 5, 6 (a) See Table 4(b) . (b) See Table 4(c) . (c) No symbol necessary; cut edge is standard. (d) Produced in five tempers with specific hardness and bend test limits; composition subordinate to mechanical properties. Source: Ref 2

Table 4(b) Selection and specification of surface condition for plain carbon steel sheet Specification symbol

Description of surface

Surface described applicable to

U(a)

Surface finish as normally used for unexposed automotive parts. Matte appearance. Normally annealed last

Cold-rolled sheet

E(b)

Surface finish as normally used for exposed automotive parts that require a Cold-rolled sheet good painted surface. Free from strain markings and fluting. Matte appearance. Temper rolled

B

Same as above, except commercial bright appearance

Cold-rolled sheet

L

Same as above, except luster appearance

Cold-rolled sheet

1

No. 1 or dull finish (no luster). Especially suitable for lacquer or paint adhesion. Facilitates drawing by reducing the contact friction between the die and the metal

Cold-rolled strip

2

No. 2 or regular bright finish (moderately smooth). Suitable for many applications, but not generally applicable for parts to be plated, unless polished and buffed

Cold-rolled strip

3

No. 3 or best bright finish (relatively high luster). Particularly suitable for parts to be plated

Cold-rolled strip

A

As-rolled or black (oxide or scale not removed)

Hot-rolled sheet and strip

P

Pickled (scale removed), not oiled


O Same as above, except oiled Hot-rolled sheet and strip (a) U, unexposed; also designated as class 2, cold-rolled sheet. (b) E, exposed; also designated as class 1, cold-rolled sheet. Source: Ref 2

Table 4(c) Selection and specification of edge condition of plain carbon steel sheet and strip Specification symbol

Description of edge

Edge described applicable to

None required

Cut edge

Cold-rolled sheet

1

No. 1 edge is a prepared edge of a specified contour (round, square, or beveled) supplied when a very accurate width is required or where the finish of the edge is required to be suitable for electroplating or both

Cold-rolled strip

2

No. 2 edge is a natural mill edge carried through the cold rolling from the hot-rolled strip without additional processing of the edge

Cold-rolled strip

3

No. 3 edge is an approximately square edge produced by slitting, on which the burr is Cold-rolled strip not eliminated

4

No. 4 edge is a rounded edge produced by edge rolling the natural edge of hot-rolled strip or slit-edge strip. This edge is produced when the width tolerance and edge condition are not as exacting as for No. 1 edge

5

No. 5 edge is an approximately square edge produced by rolling or filing of a slit edge Cold-rolled strip to remove burr only

6

No. 6 edge is a square edge produced by edge rolling the natural edge of hot-rolled strip or slit-edge strip, where the width tolerance and finish required are not as exacting as for No. 1 edge

Cold-rolled strip

M

Mill edge


C

Cut edge


S

Square edge (square and smooth, corners slightly rounded). Produced by rolling through vertical edging rolls during the hot-rolling operation

Hot-rolled strip

Cold-rolled strip

Source: Ref 2

Production of Carbon Steel Sheet and Strip Carbon steel sheet and strip are available as hot-rolled and as cold-rolled products. Hot-rolled low-carbon steel sheet and strip are usually produced on continuous hot strip mills. The slab is heated and then passed through the mill, where the thickness is progressively reduced to the desired final dimension (see Fig. 3 and the corresponding text in the section "Direct Casting


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01 Sep 2005

Methods" in this article). Some wide hot strip mills are capable of producing low-carbon steel sheet in thicknesses as low as 1.214 mm (0.0478 in.) (18 gage), but 1.897 mm (0.0747 in.) (14 gage) is considered a practical lower limit. Most narrow hot-strip mills are capable of producing low-carbon steel strip in thicknesses as low as 1.062 mm (0.0418 in.) (19 gage). Cold-rolled low-carbon steel sheet and strip are produced from pickled hot-rolled coils by cold reduction to the desired thicknesses in either a continuous tandem mill or a reversing cold-reduction mill. The cold-rolling process allows thinner gages to be produced than can be obtained by hot rolling. Other advantages of cold-rolled steel are its better surface finish and dimensional control. The as-rolled steel is hard and has low productivity. Except when a fully work-hardened condition is desired, the steel is annealed to optimize its formability. This annealing can be range from stress relieving through full recrystallization with ferrite grain growth and carbide agglomeration (see the article "Steel Processing Technology" in this Volume). After annealing, temper rolling (also called skin rolling or skin passing) is usually done to improve flatness and surface finish. Roller leveling or tension leveling can be used to improve flatness. Temper rolling, roller leveling, or tension leveling will also minimize the tendency of the material to develop stretcher strains during forming; this effect is permanent with killed steels and temporary with rimmed and capped steels (see the section "Surface Characteristics" in this article). Heating a killed steel, as in baking paint, may cause the steel to become susceptible to stretcher strains (see the article "Precoated Steel Sheet" in this Volume). Most cold-rolled low-carbon steel sheet is available in two classes (Table 4(a) ). Class 1 (temper rolled) is intended for applications where surface appearance is important and where specified surface and flatness requirements must be met. Class 2 is a product intended for applications where appearance is less important. Cold-rolled low-carbon steel strip is available in five hardness tempers ranging from full hard to dead soft (Table 5 ). Table 5 Mechanical properties of cold-rolled low-carbon steel strip (ASTM A 109) Approximate tensile strength Temper No. 1 (hard)

Hardness requirements, HRB 90 minimum(c), 84 minimum(d)

MPa

ksi

Elongatio n in 50 mm (2 in.), % (b)

No bending in either direction

550−690

80−100

...

Bent test requirements(a)

No. 2 (half-hard)

70−85(d)

90° bend across rolling direction around a 1t radius

380−520

55−75

4−16

No. 3 (quarter-hard)

60−75(e)

180° bend across rolling direction and 90° bend along rolling direction, both around a 1t radius

310−450

45−65

13−27

290−370

42−54

24−40

No. 4 (skin rolled)

65 maximum(e)

Bend flat on itself in any direction

No. 5 (dead soft)

55 maximum(e)

Bend flat on itself in any direction

260−340 38−50 33−45 (a) t = thickness of strip. (b) For strip 1.27 mm (0.050 in.) thick. (c) For strip of thickness 1.02−1.78 mm exclusive (0.040−0.070 in. exclusive). (d) For strip of thickness 1.78−6.35 mm exclusive (0.070−0.250 in. exclusive). (e) For strip of thickness 1.02−6.35 mm exclusive (0.040−0.250 in. exclusive)

Quality Descriptors for Carbon Steels The descriptors of quality used for hot-rolled plain carbon steel sheet and strip and cold-rolled plain carbon steel sheet include structural quality, commercial quality, drawing quality, and drawing quality, special killed (Table 4(a) ). Some of the as-rolled material made to these qualities is subject to surface disturbances known as coil breaks, fluting, and stretcher strains; however, fluting and stretcher strains will not be produced during subsequent forming if the material is temper rolled and/or roller leveled immediately prior to forming. It should be noted that any beneficial effects of roller leveling deteriorate rapidly in nonkilled steel. In addition to the requirements listed below for the various qualities of plain carbon steel sheet and strip, special soundness can also be specified. Commercial quality (CQ) plain carbon steel sheet and strip are suitable for moderate forming; material of this quality has sufficient ductility to be bent flat on itself in any direction in a standard room-temperature bend test. Commercial quality material is not subject to any other mechanical test requirements, and it is not expected to have exceptionally uniform chemical composition or mechanical properties. However, the hardness of cold-rolled CQ sheet is ordinarily less than 60 HRB at the time of shipment. Drawing Quality. When greater ductility or more uniform properties than those afforded by commercial quality are required, drawing quality (DQ) is specified. Drawing quality material is suitable for the production of deep-drawn parts and other parts requiring severe deformation. When the deformation is particularly severe or resistance to stretcher strains is required, drawing quality, special killed (DQSK) is specified. When either type of drawing quality material is specified, the supplier usually guarantees that the material is capable of being formed into a specified part within an established breakage allowance. The identification of the part is included in the purchase order. Ordinarily, DQ or DQSK material is not subject to any other mechanical requirements, nor is it normally ordered to a specific chemical composition. Special killed steel is usually an aluminum-killed steel, but other deoxidizers are sometimes used to obtain the desired


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characteristics. In addition to severe drawing applications, it is specified for applications requiring freedom from significant variations in mechanical properties or freedom from fluting and stretcher strains in temper-rolled material without subsequent roller leveling prior to forming. Special killed steels also have inherent characteristics that increase their formability. Structural quality (SQ), formerly called physical quality (PQ), is applicable when specified strength and elongation values are required in addition to bend tests (Table 6 ). Minimum values of tensile strength ranging up to 690 MPa (100 ksi) in hot-rolled sheet and strip and up to 1035 MPa (150 ksi) in cold-rolled sheet are available. Cold-rolled strip, which does not have a quality descriptor, is available in five tempers that conform to specified Rockwell hardness ranges and bend test requirements (Table 5 ). It should be noted that steels with yield strengths exceeding 275 MPa (40 ksi) or tensile strengths greater than 345 MPa (50 ksi) are referred to as high-strength structural or high-strength low-alloy steels. These materials are described elsewhere in this Section of the Volume (see the articles "Classification and Designation of Carbon and Low-Alloy Steels" and "High-Strength Structural and High-Strength Low-Alloy Steels" ). Table 6 Tensile requirements for hot-rolled and cold-rolled plain carbon steel sheet and strip Yield strength, minimum Class or grade

MPa

ksi

Tensile strength, minimum MPa

ksi

Elongation in 50 mm (2 in.), minimum, %

Structural quality hot-rolled sheet and strip in cut lengths or coils (ASTM A 570)(a) 30

205

30

340

49

25.0(b)

33

230

33

360

52

23.0(b)

36

250

36

365

53

22.0(b)

40

275

40

380

55

21.0(b)

45

310

45

415

60

19.0(b)

50

345

50

450

65

17.0(b)

55

380

55

480

70

15.0(b)

Structural quality cold-rolled sheet in cut lengths or coils (ASTM A 611)(a) A

170

25

290

42

26

B

205

30

310

45

24

C

230

33

330

48

22

D, types 1 and 2

275

40

360

52

20

550(c)

80(c)

565

82

...

E

Hot-rolled sheet for pressure vessels (ASTM A 414) A

170(d)

25(d)

310

45

26(e)

B

205(d)

30(d)

345

50

24(e)

C

230(d)

33(d)

380

55

22(e)

D

240(d)

35(d)

415

60

20(e)

E

260(d)

38(d)

450

65

18(e)

F

290(d)

42(d)

485

70

16(e)

G 310(d) 45(d) 515 75 16(e) (a) For coil products, testing by the producer is limited to the end of the coil. Results of such tests must comply with the specified values. However, design considerations must recognize that variation strength levels may occur throughout the untested portions of the coil, but generally these levels will not be less than 90% of the minimum values specified. (b) At thickness, t, of 2.5−5.9 mm (0.097−0.230 in.). (c) On this full-hard product, the yield point approaches the tensile strength and because there is no halt in the gage or drop in the beam, the yield point shall be taken as the stress at 0.5% elongation, under load. (d) Yield strength determined by the 0.2% offset or 0.5% extension under load methods. (e) At thickness, t, of 3.7−5.9 mm (0.145−0.230 in.). Source: Ref 1

Mechanical Properties of Carbon Steels The commonly measured tensile properties of plain carbon steel sheet and strip are not readily related to their performance in fabrication; the relationship between formability and values of the strain-hardening exponent, n, and the plastic strain ratio, r (determined in tensile testing), is discussed in the article "Sheet Formability of Steels" in this Volume. The mechanical properties of commercial quality, drawing quality, and drawing quality, special killed sheet and strip are not ordinarily used in specifications unless special strength properties are required in the fabricated product. As a matter of general interest, however, the ranges of mechanical properties typical of sheet produced by three mills in these qualities are shown in Fig. 1 . The bands would be wider if the product of the entire industry were represented. Fig. 1 Typical mechanical properties of low-carbon steel sheet shown by the range of properties in steel furnished by three


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mills. Hot-rolled sheet thickness from 1.519 to 3.416 mm (0.0598 to 0.1345 in., or 16 to 10 gage); cold-rolled sheet thickness from 0.759 to 1.519 mm (0.0299 to 0.0598 in., or 22 to 16 gage). All cold-rolled grades include a temper pass. All grades were rolled from rimmed steel except the one labeled special killed. See Table 5 for the mechanical properties of structural (physical) quality sheet.

It should be noted that the ranges are broader and the sheet harder for the hot-rolled than for the cold-rolled materials and that cold-rolled drawing quality, special killed sheet is produced to a narrower range of mechanical properties than cold-rolled drawing quality sheet, which is a rimmed steel grade. There is a great deal of overlapping in properties between commercial quality and drawing quality sheet. Figure 2 shows the relationships among hardness, Olsen ductility, and sheet thickness in commercial quality and drawing quality hot-rolled low-carbon steel sheet, indicating the variations in properties that can occur in these materials. Stretchability, as measured by the Olsen value, is also shown to increase as sheet thickness increases. Fig. 2 Scatter in Olsen ductilities of hot-rolled low-carbon steel sheet


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In contrast to commercial and drawing quality materials, structural (physical) quality sheet and strip are produced in many grades having specific mechanical property minimums, seven of which are shown in Table 6 . Cold-rolled low-carbon steel strip is not usually produced to specific strength requirements; typical mechanical property ranges for the various tempers of this product are listed in Table 5 . Mill Heat Treatment of Cold-Rolled Products Unless a hard temper is desired, cold-rolled carbon steel sheet and strip are always softened to improve formability. This is usually accomplished at the mill by a recrystallization heat treatment such as annealing or normalizing. Annealing. Low-temperature recrystallization annealing, or process annealing, can be used to soften cold-rolled low-carbon steel. When done as a batch process, this type of annealing is known as box annealing. It is carried out by placing coils on a bottom plate and then enclosing them with a cover within which a protective gas atmosphere is maintained. A bell-type heating furnace is then placed over the atmosphere container. After heating to approximately 595 to 760 °C (1100 to 1400 °F), the charge is allowed to soak until the temperature is uniform throughout. The heating furnace is then removed, and the charge is allowed to cool in the protective atmosphere before being uncovered. Cold-rolled steel can be batch annealed in coil form under a protective atmosphere. Some producers use a 100% hydrogen atmosphere in an effort to shorten annealing cycles. Instead of box annealing, coils can also be treated by continuous annealing. With this process, which is usually intended to provide a fully recrystallized grain structure, coils are unwound and passed through an annealing furnace. The uncoiled steel strip passes through several different thermal zones of the furnace that serve to heat, soak, and cool the steel before it exits the furnace and is recoiled. This anneal cycle is very rapid and can be measured in seconds or minutes (as opposed to hours or days with a box anneal cycle). Generally, the rapid anneal cycle of a continuous anneal process results in material properties that are less ductile than those resulting from a box anneal cycle. However, continuous annealing results in more uniformity of properties throughout the length of a coil. Open-coil annealing is used when uniform heating and/or gas contact across the entire width of the coil is required (for example, to obtain decarburization over the entire surface during production of material for porcelain enameling). In this process, the coils are loosely wound, permitting gas to flow freely between the coil convolutions. Annealing temperatures may be higher than those used in conventional box annealing. Normalizing consists of heating the sheet or strip to a temperature above the Ac3 point (~925 °C, or 1700 °F, for a steel that contains less than 0.15% C) in a continuous furnace containing an oxidizing atmosphere, then cooling to room temperature at a controlled rate (usually in still air). This treatment recrystallizes and refines the grain structure by phase transformation. Low-metalloid steel (enameling iron) for porcelain enameling is normalized rather than annealed because this steel will not readily recrystallize at box-annealing temperatures. Surface Characteristics The surface texture of low-carbon cold-rolled steel sheet and strip can be varied between rather wide limits. For chromium plating and similar finishes, a smooth, bright sheet or strip surface is necessary, but for porcelain enameling and many drawing operations, a rougher surface texture (matte finish) is preferred. In porcelain enameling, roughness tends to improve the adherence and uniformity of the coating; in certain drawing operations where heavy pressures are developed, the rougher type of surface is believed to retain more lubricant, thus aiding formation of the sheet by reducing friction and die galling. Minor surface imperfections and slight strains are less noticeable on a dull surface than on a bright one. However, the surfaces of parts to be painted should not be so rough that the paint will not cover them adequately. A very smooth, bright surface can be obtained on sheet or strip by utilizing ground and polished rolling-mill rolls, and a dull (matte) surface can be obtained by either grit blasting or etching the rolls. For the purpose of evaluating surface roughness, an appropriate instrument is employed that measures the average height of surface asperities (peaks) in microinches and the number of peaks per inch that exceed a given height. Cold-rolled sheet or strip can also be purchased with coined patterns that form a geometric design or that simulate such textures as leather grain. Such products are available in commercial quality, drawing quality, and drawing quality, special killed material. The texture is rolled into the steel surface after the sheet or strip has been annealed and thus has an effect on properties similar to that of a heavy temper-rolling pass. This effect, plus the notch effect of the pattern itself, somewhat reduces the formability of the sheet or strip. Stretcher Strains. When loaded in tension, practically all hot-rolled or as-annealed cold-rolled plain carbon steels, whether rimmed, capped, or killed, exhibit a sharp upper yield point, a drop in load to the lower yield point, and subsequent plastic deformation at a nearly constant load (known as yield point elongation). The plastic deformation that occurs within this yield point elongation is accompanied by the formation of visible bands of deformation on the product surfaces. These bands are called stretcher strains or Lüders lines, and they can be aesthetically undesirable. The tendency for stretcher strains to occur can be prevented through elimination of yield point elongation. In rimmed or capped steels, this is accomplished by subjecting the steel to small amounts of plastic deformation, usually by temper rolling, tension leveling, and/or roller leveling. Because overstraining the steel by these practices can increase strength and generally decrease ductility, it is usually desirable to strain the steel only by the amount required to eliminate yield point elongation. When properly processed, a killed steel, such as DQSK, provides a product with no yield point elongation.


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Strain Aging. In rimmed or capped (but not killed) carbon steels, deformation (such as by temper rolling) following by aging for several days or more at or slightly above room temperature will result in a return of the upper yield point and yield point elongation, increases in yield and tensile strengths, and a decrease in ductility. This treatment, called strain aging, may be desirable if the increase in strength can be used to advantage. However, strain aging often causes problems due to reduced formability and stretchability and the return of both yield point elongation and a propensity for stretcher strains. Further temper rolling may eliminate yield point elongation, but it will not restore stretchability. In applications where the appearance of stretcher strains is objectionable, killed steels, which are resistant to aging, are preferable to rimmed and capped steels. For ingot casting, however, rimmed and capped steels are generally superior in inherent surface quality, are lower in cost, and are preferred over killed steel as long as the occurrence of stretcher strains is not a problem. Strain aging is related to the presence of nitrogen in solid solution in the steel and is affected by time and temperature, with longer times and higher temperatures producing greater aging. The strain-aging rate is also dependent on the amount of deformation that has occurred and is increased when the deformation occurs at higher temperatures or lower strain rates. Another important variable that affects strain aging is the amount of nitrogen in solution. Killed carbon steels have very little susceptibility to strain aging because their nitrogen content is essentially chemically combined with aluminum. Rimmed and capped steels, however, tend to strain age because they contain greater amounts of nitrogen in solid solution (typically 6 to 30 ppm). Control of Flatness Plain carbon steel sheet is ordinarily sold to two standards of flatness: • Commercial flatness, which is used where flatness is important but not critical • The stretcher-level standard of flatness, which is required when little or no forming is to be done and the product is required to be flat and free from waves or oil can, or when flatness is necessary to ensure smooth automatic feeding of forming equipment. The permissible variations for the flatness of hot- and cold-rolled sheet have been established by the Technical Committee of the American Iron and Steel Institute and are given in the AISI Steel Products Manual. Commercial flatness can usually be produced by roller leveling or by temper rolling and roller leveling, but where very flat sheet is required, producers may have to resort to stretcher leveling, tension leveling, or other leveling processes. In temper rolling, the steel is cold reduced, usually by 1=2to 2%, which is also effective for removing yield point elongation and preventing stretcher strains. In roller leveling, a staggered series of small-diameter rolls alternately flexes the steel back and forth. The rolls are adjusted so that the greatest deformation occurs at the entrance end of the rolls and less flexing occurs at the exit end. Stretcher strains can also be eliminated by roller leveling, as long as the deformation is great enough to remove yield point elongation. Dead-soft annealed sheet cannot be made suitable for production of exposed parts by roller leveling because the rolls kink the sheet severely, producing leveler breaks. The deformed areas or kinks will not deform further upon stretching and will appear as braised welts after forming. Stretcher Leveling. Leveling by stretching cut lengths of the temper-rolled sheet lengthwise between jaws (stretcher leveling) is a more positive means of producing flatness. Elongation (stretching) during stretcher leveling may vary from about 1 to 3%, which exceeds the elastic limit of the steel and therefore results in some permanent elongation. The sheet must be of a killed or a capped steel having nearly uniform properties so that it will spring back uniformly across its full width and remain flat. It may be necessary to use killed steel having nearly uniform properties so that, after stretching, strain markings do not develop. Tension Leveling. Another flattening process that is used for steel sheet is tension leveling, which combines the effects of stretcher and roller leveling. The sheet is pulled to a stress near its yield point while it is simultaneously flexed over small rolls; the combined tension and bending produce yielding at the flex points. Modified Low-Carbon Steel Sheet and Strip In addition to the low-carbon steel sheet and strip products already discussed in this article, there are numerous additional products available that are designed to satisfy specific customer requirements. These products are often made with low-carbon steels having chemical compositions slightly modified from those discussed earlier. To be considered a plain low-carbon grade, a steel should contain no more than 0.25% C, 1.65% Mn, 0.60% S, and 0.60% Cu, but it may also contain small amounts of other elements, such as nitrogen, phosphorus, and boron, that are effective in imparting special characteristics when present singly or in combination. The modified low-carbon steel grades discussed below are designed to provide sheet and strip products having increased strength, formability, and/or corrosion resistance. Carbon-Manganese Steels. Manganese is a solid-solution strengthening element in ferrite and is also effective in increasing hardenability. Manganese in amounts ranging from 1.0 to 1.5% is added to low-carbon steel (0.15 to 0.25% C) to provide enhanced strength (yield strength of about 275 MPa, or 40 ksi) with good ductility in hot-rolled and cold-rolled sheet and strip. Components fabricated from these higher-manganese steels can be heat treated by quenching and tempering to provide enhanced strength with good toughness (see the article "High-Strength Structural and High-Strength Low-Alloy Steels" in this Volume).


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Carbon-Silicon Steels. Silicon, like manganese, is an effective ferrite-strengthening element and is sometimes added in amounts of about 0.5%, often in combination with 1.0 to 1.5% Mn, to provide increased strength in low-carbon hot-rolled and cold-rolled steel sheet and strip. Nitrogenized and Rephosphorized Steels. Nitrogen is a strong interstitial strengthener, and phosphorus is an effective solid-solution strengthener in ferrite. Either about 0.010 to 0.015% N or 0.07 to 0.12% P is added to low-carbon steel to provide hot-rolled and cold-rolled sheet and strip with yield strength in the range of 275 to 345 MPa (40 to 50 ksi) for low-cost structural components for buildings and automotive uses. Formed parts produced from nitrogenized steel can be further strengthened to yield strengths in the range of 415 to 485 MPa (60 to 70 ksi) as the result of strain aging that occurs at paint-curing temperatures. Boron Steels. Boron is a strong carbide-and nitride-forming element and increases strength in quenched and tempered low-carbon steels through the formation of martensite and the precipitation strengthening of ferrite. Boron-containing killed carbon steels are available as low-cost replacements for the high-carbon and low-alloy steels used for sheet and strip. The low-carbon boron steels have better cold-forming characteristics and can be heat treated to equivalent hardness and greater toughness for a wide variety of applications, such as tools, machine components, and fasteners. Copper Steels. Copper in amounts up to 0.5% is not only a mild solid-solution strengthener in ferrite, but it also provides enhanced atmospheric corrosion resistance together with improved paint retention in applications involving full exposure to the weather. Therefore, copper-bearing (0.20% Cu, minimum) steel is often specified by customers for use in sheet and strip for structures subject to atmospheric corrosion. Essentially all low-carbon steel sheet and strip products can be supplied in copper-bearing grades, if so specified. Copper-bearing steels, which are also referred to as weathering steels, are also described in the article "High-Strength Structural and High-Strength Low-Alloy Steels" in this Volume.

Low-Alloy Steel* [*The term low-alloy steel rather than the more general term alloy steel is being used in this article as well as other articles in this Section of the Handbook. See the article "Classification and Designations of Carbon and Low-Alloy Steels" for definitions of various steel types.] Low-alloy steel sheet and strip are used primarily for those special applications that require the mechanical properties normally obtained by heat treatment. A sizeable selection of the standard low-alloy steels are available as sheet and strip, either hot-rolled or cold rolled. The most commonly available alloys are listed in Table 7 , along with their chemical compositions. In addition to standard low-alloy steels, high-strength low-alloy (HSLA) and dual-phase steels are available as sheet or strip for applications requiring tensile strengths in the range of 290 to 760 MPa (42 to 110 ksi), and ultrahigh-strength steels or maraging steels for applications requiring tensile strengths above 1380 MPa (200 ksi). These steels are discussed in the articles "High-Strength Structural and High-Strength Low-Alloy Steels," "Dual-Phase Steels," "Ultrahigh-Strength Steels" and "Maraging Steels" in this Volume. Table 7 Compositions for hot-rolled and cold-rolled low-alloy steel sheet and strip AISI or SAE designation

Chemical composition ranges and limits, % (heat analysis)(a) C

Mn

P

S

Si(b)

Ni

Cr

Mo

V

Regular quality and structural quality standard steels commonly produced (ASTM A 506) 4118

0.18−0.23

0.70−0.90

0.035

0.040

0.15−0.30

...

0.40−0.60

0.08−0.15

...

4130

0.28−0.33

0.40−0.60

0.035

0.040

0.15−0.30

...

0.80−1.10

0.15−0.25

...

4140

0.38−0.43

0.75−1.00

0.035

0.040

0.15−0.30

...

0.80−1.10

0.15−0.25

...

4340

0.38−0.43

0.60−0.80

0.035

0.040

0.15−0.30

1.65−2.00

0.70−0.90

0.20−0.30

...

0.040

0.15−0.30

...

0.70−0.90

...

...

5140

0.38−0.43

0.70−0.90

0.035

5150

0.48−0.53

0.70−0.90

0.035

0.040

0.15−0.30

...

0.70−0.90

...

...

5160

0.55−0.65

0.75−1.00

0.035

0.040

0.15−0.30

...

0.70−0.90

...

...

8615

0.13−0.18

0.70−0.90

0.035

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

8620

0.18−0.23

0.70−0.90

0.035

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

... ...

Regular quality and structural quality standard steels not commonly produced (ASTM A 506) E3310

0.08−0.13

0.45−0.60

0.025

0.025

0.15−0.30

3.25−3.75

1.40−1.75

...

4012

0.09−0.14

0.75−1.00

0.040

0.040

0.15−0.30

...

...

0.15−0.25

...

0.70−0.90

0.040

0.040

0.15−0.30

...

0.40−0.60

0.08−0.15

...

0.040

0.15−0.30

...

0.80−1.10

0.15−0.25

...

4118

0.18−0.23

4135

0.33−0.38

0.70−0.90

0.040

4137

0.35−0.40

0.70−0.90

0.040

0.040

0.15−0.30

...

0.80−1.10

0.15−0.25

...

4142

0.40−0.45

0.75−1.00

0.040

0.040

0.15−0.30

...

0.80−1.10

0.15−0.25

...

4145

0.43−0.48

0.75−1.00

0.040

0.040

0.15−0.30

...

0.80−1.10

0.15−0.25

...


Page 346



01 Sep 2005

4147

0.45−0.50

0.75−1.00

0.040

0.040

0.15−0.30

...

0.80−1.10

0.15−0.25

...

4150

0.48−0.53

0.75−1.00

0.040

0.040

0.15−0.30

...

0.80−1.10

0.15−0.25

...

4320

0.17−0.22

0.45−0.65

0.040

0.040

0.15−0.30

1.65−2.00

0.40−0.60

0.20−0.30

...

E4340

0.38−0.43

0.65−0.85

0.025

0.025

0.15−0.30

1.65−2.00

0.70−0.90

0.20−0.30

...

4520

0.18−0.23

0.45−0.65

0.040

0.040

0.15−0.30

...

...

0.45−0.60

...

0.040

0.15−0.30

1.65−2.00

...

0.20−0.30

...

4615

0.13−0.18

0.45−0.65

0.040

4620

0.17−0.22

0.45−0.65

0.040

0.040

0.15−0.30

1.65−2.00

...

0.20−0.30

...

4718

0.16−0.21

0.70−0.90

0.040

0.040

0.15−0.30

0.90−1.20

0.35−0.55

0.30−0.40

...

4815

0.13−0.18

0.40−0.60

0.040

0.040

0.15−0.30

3.25−3.75

...

0.20−0.30

...

0.040

0.15−0.30

3.25−3.75

...

0.20−0.30

...

4820

0.18−0.23

0.50−0.70

0.040

5015

0.12−0.17

0.30−0.50

0.040

0.040

0.15−0.30

...

0.30−0.50

...

...

5046

0.43−0.50

0.75−1.00

0.040

0.040

0.15−0.30

...

0.20−0.35

...

...

5115

0.13−0.18

0.70−0.90

0.040

0.040

0.15−0.30

...

0.70−0.90

...

...

5130

0.28−0.33

0.70−0.90

0.040

0.040

0.15−0.30

...

0.80−1.10

...

...

0.040

0.15−0.30

...

0.75−1.00

...

...

5132

0.30−0.35

0.60−0.90

0.040

E51100

0.95−1.10

0.25−0.45

0.025

0.025

0.15−0.30

...

0.90−1.15

...

...

E52100

0.95−1.10

0.25−0.45

0.025

0.025

0.15−0.30

...

1.30−1.60

...

...

6150

0.48−0.53

0.70−0.90

0.040

0.040

0.15−0.30

...

0.80−1.10

...

0.15 min

8617

0.15−0.20

0.70−0.90

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

8630

0.28−0.33

0.70−0.90

0.040

8640

0.38−0.43

0.75−1.00

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

8642

0.40−0.45

0.75−1.00

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

8645

0.43−0.48

0.75−1.00

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

8650

0.48−0.53

0.75−1.00

0.040

8655

0.50−0.60

0.75−1.00

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

8660

0.55−0.65

0.75−1.00

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

8720

0.18−0.23

0.70−0.90

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.20−0.30

...

8735

0.33−0.38

0.75−1.00

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.20−0.30

...

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.20−0.30

...

8740

0.38−0.43

0.75−1.00

0.040

9260

0.55−0.65

0.70−1.00

0.040

0.040

1.80−2.20

...

...

...

...

9262

0.55−0.65

0.75−1.00

0.040

0.040

1.80−2.20

...

0.25−0.40

...

...

E9310

0.08−0.13

0.45−0.65

0.025

0.025

0.20−0.35

3.00−35.0

1.00−1.40

0.08−0.15

...

Drawing quality standard steels commonly produced (ASTM A 507) 4118

0.18−0.23

0.70−0.90

0.035

0.040

0.15−0.30

...

0.40−0.60

0.08−0.15

...

4130

0.28−0.33

0.40−0.60

0.035

0.040

0.15−0.30

...

0.80−1.10

0.15−0.25

...

8615

0.13−0.18

0.70−0.90

0.035

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

0.70−0.90

0.035

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

8620

0.18−0.23

Drawing quality standard steels not commonly produced (ASTM M A 507) E3310

0.45−0.60

0.025

0.025

0.15−0.30

3.25−3.75

1.40−1.75

...

...

0.09−0.14

0.75−1.00

0.040

0.040

0.15−0.30

...

...

0.15−0.25

...

4118

0.18−0.23

0.70−0.90

0.040

0.040

0.15−0.30

...

0.40−0.60

0.08−0.15

...

4320

0.17−0.22

0.45−0.65

0.040

0.040

0.15−0.30

1.65−2.00

0.40−0.60

0.20−0.30

...

4520

0.18−0.23

0.45−0.65

0.040

0.040

0.15−0.30

...

...

0.45−0.60

...

4615

0.13−0.18

0.45−0.65

0.040

0.040

0.15−0.30

1.65−2.00

...

0.20−0.30

...

0.040

0.15−0.30

1.65−2.00

...

0.20−0.30

...

0.040

0.15−0.30

0.90−1.20

0.35−0.55

0.30−0.40

...

4012

0.08−0.13

4620

0.17−0.22

0.45−0.65

0.040

4718

0.16−0.21

0.70−0.90

0.040


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01 Sep 2005

4815

0.13−0.18

0.40−0.60

0.040

0.040

0.15−0.30

3.25−3.75

...

0.20−0.30

...

4820

0.18−0.23

0.50−0.70

0.040

0.040

0.15−0.30

3.25−3.75

...

0.20−0.30

...

5015

0.12−0.17

0.30−0.50

0.040

0.040

0.15−0.30

...

0.30−0.50

...

...

5115

0.13−0.18

0.70−0.90

0.040

0.040

0.15−0.30

...

0.70−0.90

...

...

5130

0.28−0.33

0.70−0.90

0.040

0.040

0.15−0.30

...

0.80−1.10

...

...

0.040

0.15−0.30

...

0.75−1.00

...

...

5132

0.30−0.35

0.60−0.90

0.040

8617

0.15−0.20

0.70−0.90

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

8630

0.28−0.33

0.70−0.90

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.15−0.25

...

8720

0.18−0.23

0.70−0.90

0.040

0.040

0.15−0.30

0.40−0.70

0.40−0.60

0.20−0.30

...

E9310 0.025 0.025 ... 0.08−0.13 0.45−0.65 0.20−0.35 3.00−3.50 1.00−1.40 0.08−0.15 (a) The chemical ranges and limits shown are subject to product analysis tolerances. See ASTM A 505. (b) Other silicon ranges are available. Consult the producer. Source: Ref 1

Production of Sheet and Strip As described earlier in this article, steel sheet and strip are flat-rolled products that can be rolled to finished thickness on either a hot mill or a cold mill. Hot-rolled steel sheet and strip are normally produced by passing heated slabs through a continuous mill consisting of a series of roll stands, where the thickness is progressively reduced to the desired final dimension. Cold-rolled low-alloy steel sheet and strip are normally produced from pickled and annealed hot-rolled bands of intermediate thickness by cold reduction to desired thickness in a single-stand mill or tandem mill. Intermediate anneals may be required to facilitate cold reduction or to obtain the mechanical properties desired in the finished product. Cold rolling can produce thinner gages than can be obtained by hot rolling. Low-alloy steel sheet and strip are produced in thicknesses similar to those typical of HSLA steel sheet and strip (Table 8 ). In general, tolerances similar to those given in the general requirements for hot-rolled and cold-rolled low-alloy and HSLA steel sheet and strip, ASTM A 505, apply to all low-alloy and HSLA steel sheet and strip. Available thicknesses and tolerances may vary among producers, due mainly to the interrelation between steel quality and rolling practice, as influenced by the equipment available for rolling the product. Table 8 Standard sizes of hot-rolled and cold-rolled low-alloy steel sheet and strip: regular quality, structural quality, and drawing quality Product Hot-rolled sheet

Hot-rolled strip

Cold-rolled sheet Cold-rolled strip Source: Ref 1

Applicable ASTM specification A 506, A 507

A 506, A 507

A 506, A 507 A 506, A 507

Thickness range

Width range

mm

in.

mm

in.

5.839−4.572 inclusive

0.2299−0.1800 inclusive

610−1220 inclusive

24−48 inclusive

≤4.569

≤0.1799

>610

>24

≤5.156

≤0.2030

≤152

≤5.839

≤0.2299

152−608 inclusive

>6−23 =16 inclusive

≤6 15

≤5.839

≤0.2299

610−1220

24−48

≤4.569

≤0.1799

>1220

>48

≤6.347

≤0.2499

≤608

≤2315=16

Quality Descriptors As it is used for steel mill products, the term quality relates to the general suitability of the mill product to make a given class of parts. For low-alloy steel sheet and strip, the various quality descriptors imply certain inherent characteristics, such as the degree of internal soundness and the relative freedom from harmful surface imperfections. The quality descriptors used for alloy steel sheet and plate include regular quality, drawing quality, and aircraft quality, which are covered by ASTM specifications. The general requirements for these qualities include bearing quality and aircraft structural quality. Aircraft quality requirements are also defined in Aerospace Material Specifications (AMS). Regular Quality. Low-alloy steel sheet and strip of regular quality are intended principally for general or miscellaneous applications where moderate drawing and/or bending is required. A smooth finish free of minor surface imperfections is not a primary requirement. Sheet and strip of this quality do not have the uniformity, the high degree of internal soundness, or the freedom from surface imperfections that are associated with other quality descriptors for low-alloy sheet and strip. Regular quality low-alloy steel sheet and strip are covered by ASTM A 506. One or more of the following characteristics may


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01 Sep 2005

be specified by the purchaser: chemical composition, grain size, or mechanical properties (determined by tensile and bend tests.) Drawing quality describes low-alloy steel sheet and strip for applications involving severe cold working such as deep-drawn or severely formed parts. Drawing quality low-alloy sheet and strip are rolled from steel produced by closely controlled steelmaking practices. The semifinished and finished mill products are subject to testing and inspection designed to ensure internal soundness, relative uniformity of chemical composition, and freedom from injurious surface imperfections. Spheroidize annealing is generally specified so the mechanical properties and microstructure are suitable for deep drawing or severe forming. Drawing quality low-alloy steel sheet and strip are covered by ASTM A 507. No standard test can fully evaluate resistance to breakage during deep drawing because successful drawing is affected by die clearances, die design, speed of drawing, lubricants, ironing, grade of steel, and any alteration of hardness, ductility, or surface condition that may develop during drawing. Thus, it cannot be assumed that merely specifying drawing quality steel will ensure a capability for drawing or forming a specific part under a given set of manufacturing conditions. Manufacturing trials may be necessary before purchase orders can be written for production material. Bearing quality describes low-alloy steel sheet and strip intended for antifriction bearing parts. The steels are generally AISI-SAE alloy carburizing grades or AISI-SAE high-carbon chromium grades. These steels are produced using steelmaking and conditioning practices that are intended to optimize internal soundness and to provide a known size, shape, and distribution of non-metallic inclusions. Standards of acceptance for microstructural quality are commonly reviewed and agreed upon between producer and purchaser for each order. Alternatively, internal soundness and microcleanliness can be determined by using immersion ultrasonic testing techniques to agreed-upon acceptance standards. More detailed information on low-alloy bearing steels can be found in the article "Bearing Steels" in this Volume. Aircraft quality describes low-alloy steel sheet and strip for important or highly stressed parts of aircraft, missiles, and similar applications involving stringent performance requirements, especially in terms of internal cleanliness. The special mill practices required for producing aircraft quality sheet and strip include careful selection of the raw materials charged into the melting furnace, exceptionally close control of the steelmaking process, cropping and discarding more of the ingot than is normal during primary reduction, selection of specific heats or portions of heats for fulfillment of a given customer order, and using exceptionally close control over process variables during reheating and rolling. Aircraft quality low-alloy steel sheet and strip generally have an austenitic grain size predominantly ASTM No. 5 or finer, with grains as coarse as ASTM No. 3 permissible. Grain size tests are normally made on rerolling slabs or billets. Aircraft quality low-alloy steel sheet and strip are covered by Aerospace Material Specifications (AMS 6454A, for example). Material of this quality is ordinarily certified that it has been produced as aircraft quality. Aircraft structural quality low-alloy steel sheet and strip meet all the requirements of aircraft quality mill products described above. In addition, they meet specified requirements for mechanical properties, which may include tensile strength, yield strength, elongation, bend test results, or results of other similar tests. Many specimens from each heat must be tested to ensure compliance with the required mechanical properties. Mill Heat Treatment Hot-rolled regular quality low-alloy steel sheet and strip are normally available from the producer either as-rolled or heat treated. Standard mill heat-treated conditions are annealed, normalized, or normalized and tempered. Cold-rolled regular quality product is normally available only in the annealed condition. Hot-rolled and cold-rolled drawing quality alloy steel sheet and strip are normally furnished by the producer in the spheroidize-annealed condition. They can be purchased in the as-rolled condition if they are to be spheroidize annealed by the user. Aircraft quality products are normally furnished in a heat-treated condition. Hot-rolled products may be annealed, spheroidize annealed, normalized, or normalized and tempered by the producer. Cold-rolled products are normally furnished only in the annealed or spheroidize-annealed condition. Annealing is done by heating the steel to a temperature near or below the lower critical temperature and holding at that temperature for a sufficient period, followed by slow cooling in the furnace. This process softens the sheet or strip for further processing, but not to the same degree as spheroidize annealing. Spheroidize annealing involves prolonged heating at a temperature near or slightly below the lower critical temperature, followed by slow cooling. The objective of this process is to change the form of the carbides in the microstructure to a globular (spheroidal) shape, which produces the greatest degree of softening. Normalizing consists of heating the sheet or strip to a temperature 55 to 70 °C (100 to 125 °F) above Ac3 and then cooling to room temperature at a controlled rate (usually in still air). This treatment recrystallizes and refines the grains by phase transformation and can be used to obtain the desired mechanical properties. Tempering consists of reheating steel to a predetermined temperature below the lower critical temperature, holding for a specified length of time, and then cooling under suitable conditions. When it is carried out as part of a mill heat treatment, tempering is done after normalizing to obtain the desired mechanical properties by modifying the as-normalized microstructure. Quenching and tempering (or hardening) is normally reserved for the user to apply as one of the final steps in the fabricating process. Mechanical Properties


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01 Sep 2005

In most instances, the mechanical properties of low-alloy steel furnished by the producer are of little consequence because they will be altered by heat treatment during fabrication. For low-alloy steel sheet and strip to be used in the mill condition, mechanical properties will vary, depending on both chemical composition and mill processing. Table 9 lists typical tensile properties for chromium-molybdenum low-alloy steel sheet and strip used for pressure vessels. Usually, low-alloy steel sheet and strip are custom produced to fulfill specific customer orders. Where necessary, any mechanical property requirements can be made part of the purchase order. Table 9 Tensile requirements of chromium-molybdenum alloy steel sheet and strip for pressure vessels (ASTM A 873) Yield strength, minimum

Tensile strength, minimum

Class

MPa

ksi

MPa

ksi

t=3.8−5.9 mm (0.145−0.230 in.)

t=1.8−3.7 mm (0.070−0.144 in.)

1

205

30

415

60

15

12

2

310

45

515

75

13

10

3

415

60

585

85

12

9

4

515

75

655

95

11

8

690

100

895

130

7

4

5 Source: Ref 1

Elongation in 50 mm (2 in.), minimum, at thickness t, %

Because the chief benefits of low-alloy steel sheet and strip accrue to the user only after the finished part is heat treated, the mechanical properties of heat-treated low-alloy steels are the ones of greatest importance. These properties can be determined from hardenability curves (see the article "Hardenability Curves" in this Volume) and heat-treating guides such as those found in the articles "Hardenable Carbon and Low-Alloy Steels" and "Hardenability of Carbon and Low-Alloy Steels" in this Volume. In general, only those properties typical of through-hardened steel of the specific grade under consideration need to be considered. Except for the most shallow hardening grades used at thicknesses at or near the upper limit for sheet and strip, parts made of low-alloy steel sheet or strip will through harden when quenched. Many grades will through harden when quenched in a slow medium such as oil and may even through harden when air cooled. The possibility of oil quenching or air cooling should always be considered for hardening thin parts, especially when warping or distortion during hardening need to be minimized. Parts made of low-alloy steel sheet and strip are sometimes carburized or carbonitrided to improve the mechanical properties or wear resistance of the surface layer. In some cases, parts that are difficult to form when made of a medium-carbon low-alloy steel can be formed from low-carbon low-alloy steel and then carburized to a uniform but higher carbon content.

Direct Casting Methods Because of the large investment needed to build conventional steelmaking casting and rolling facilities, the focus over the last ten years has been on reducing production costs and simplifying the overall steelmaking process. For the most part, cost savings have been achieved by the progression of casting technology from ingot to continuous casting, which eliminates soaking and breakdown hot rolling of large ingots. The following table compares the continuous cast share (in percent) for the United States, the European Economic Community (EEC), Japan, and the total world: Country United States

1981, %

1989, %

1990, %

20.3

63.7

66.2

EEC

42.5

73.7

74.6

Japan

70.7

94.6

95.1

24.3

44.4

46.8

Total world Source: Wharton Econometric Forecasting Associates

Conventional continuous casting of steels requires the casting of a 150 to 250 mm (6 to 10 in.) thick by 800 to 2200 mm (31 to 86 in.) wide slab that is subsequently rolled down to a thickness of 1.5 to 25 mm (0.05 to 1.0 in.) utilizing a hot strip mill having both four-stand roughing and six- or seven-stand finishing mills (Fig. 3 ). This process requires a high degree of reduction and the equivalent input of energy. Fig. 3 Key components of a continuous casting operation. Source: SMS Engineering, Inc.


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01 Sep 2005

Direct casting processes are alternatives to conventional slab casting processes. Direct casting processes for steel flat products could be defined as any casting process that produces a casting as close as possible to the final product dimensions of the next processing step. By this definition, direct casting could also be termed near-net shape casting because the final cast dimensions would approach the final product dimensions (Ref 3). Presently, there are three direct casting alternatives. Listed in increasing order according to how close they come to producing near-net shape dimensions, these processes are (Ref 3): • Thin slab casting • Thin strip casting • Spray casting The flowcharts in Fig. 4 summarize the key operations involved in these three alternative direct casting processes and compare them with those of a continuous casting process in an integrated steel production facility. Fig. 4 Flowchart of operations for various strip casting processes. Source: Ref 3


Page 351



01 Sep 2005

Thin Slab Casting. Of the three direct casting processes listed above, only the thin slab casting process is being used commercially. In thin slab casting, a slab 40 to 60 mm (1.5 to 2.5 in.) is produced. Hot rolling is not completely eliminated in this process, but the amount of reduction necessary to produce strip is greatly reduced. However, the need for a heating furnace and a roughing mill is eliminated (Fig. 5 ). In addition, thin slab casting yields a finer grain structure and a better finish than that obtained with conventional continuous casting technology. Fig. 5 Key components of a thin slab casting facility. Compare with Fig. 3 . Source: SMS Engineering, Inc.

Table 10 lists some of the countries and specific firms engaged in research and development of thin slab casting worldwide. References 3, 4, 5, 6, and 7 provide detailed information on the start-up of a thin slab casting minimill. Table 10 Alternative sheet, strip, and slab casting techniques Country

Company

Caster type

Thin slab casting United States

Bethlehem-USX Nucor

Hazelett Hazlett SMS-Concast

Great Britain

British Steel

Travelling block mold

Germany

SMS-Concast

Vertical static mold

Krupp

Hazelett

Mannesmann

Vertical static mold

Japan

Kawasaki Steel

Vertical twin belt Horizontal twin belt

Sumitomo Metals

Hazelett

Hitachi-Korf

Wheel and belt

Nippon Steel

Twin belt

Switzerland

Alusuisse

Twin block mold

Austria

Hitachi-Korf

Wheel and belt

BSC-Armco-Inland-Weirton

Twin roll

Thin strip casting United States


Page 352


Japan

France Italy


Armco

Single roll

Allegheny-Ludlum

Single roll

Argonne National Labs

Electromagnetic levitation

United Technologies

Single roll

LTV

Drum in drum

Nippon Steel

Twin roll

Kawasaki Steel

Twin roll

Nippon Kokan

Twin roll

Nippon Metals

Twin roll

Kobe

Twin roll

Nippon Yakin

Twin roll

IRSID

Twin roll

CMS

Twin roll

Danieli

Thin strip

Austria

Voest-Alpine

Single roll

Switzerland

Concast

Single roll

Germany

Mannesman-Battelle

Single substrate

01 Sep 2005

Twin roll

Spray casting Great Britain

Osprey Metals Ltd

Osprey process

Sprayforming Developments

Spray forming

Aurora Metals

Controlled spray deposition

University of Swansea

Spray forming

Sweden

Sandviken

Osprey process

Germany

Mannesmann Demag

Spray forming sheet by Osprey process

Japan

Sumitomo

Osprey process

United States

M.I.T.

Dynamic liquid compaction

Drexel University

Osprey process

Source: Ref 3

In thin strip casting, a strip that is generally less than 5 mm (0.2 in.) thick is cast. In this process, the most optimistic scenario is that the need for a hot strip mill will be eliminated altogether. As indicated in Table 10 , there are three areas of concentration in thin strip casting: • Single-roll process • Twin-roll process • Electromagnetic levitation Strip casting is expected to be available for commercialization within the next five to ten years if significant advances in control and quality can be achieved. Currently, single-roll casting is closer to commercialization processes than twin-roll processes, especially in the area of stainless steel manufacture. Additional information on thin strip casting can be found in Ref 3. In spray forming, a liquid metal is atomized and sprayed onto a substrate in an inert atmosphere to form a sheet (Ref 3). Because it eliminates conventional casting and hot rolling processes, spray forming is a true near-net shape casting technology. Compaction after forming is normally necessary to eliminate porosity and to achieve high density. This technology has been applied to the manufacture of rings, tubes, small billets, and pipes for both ferrous and non-ferrous applications. Both centrifugal atomization processes such as controlled spray deposition and gas atomization processes are included in this category (Table 10 ). The commercialization of spray casting for strip production is at least five to ten years in the future for bulk steelmaking. In addition, applying this technology to low-carbon aluminum-killed strip may be difficult because of surface quality and yield requirements. REFERENCES 1. Steel⎯Plate, Sheet, Strip, Wire, Vol 01.03. Annual Book of ASTM Standards, American Society for Testing and Materials 2. Materials, Vol 1, SAE Handbook, Society of Automotive Engineers, 1989 3. A.W. Cramb, New Steel Casting Processes for Thin Slabs and Strip: A Historical Perspective, Iron Steelmaker, Vol 15 (No.


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01 Sep 2005

7), July 1988 4. W.D. Huskonen, Nucor Starts Up Thin Slab Mill, 33 Met. Prod., Aug 1989 5. G.J. McManus, Taking the Wraps off Nucor's Sheet Mill, Iron Age, June 1989 6. G. Flemming, F. Hollmann, M. Kolakowski, and H. Streubel, Continuous Casting of Strips, CSP: A Future Alternative for the Modernization of Slab Production, Fachber. Hüttenprax. Metallweiterverarb., Vol 25 (No. 8), 1987 7. A Collier, Hot Tech: Thin Slabs and Direct Steelmaking, Iron Age, July 1989


Page 354


Precoated Steel Sheet

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Carbon and Low-Alloy Steels Precoated Steel Sheet Revised by R. W. Leonard, USS Corporation, Division of USX Corporation STEEL SHEET is often coated in coil form before fabrication either by the steel mills or by specialists known as coil coaters. This prefinished or precoated sheet is ready for fabrication and use without further surface coating. Precoated products yield lower production costs, improved product quality, shorter processing cycles, elimination of production hazards, conservation of energy, minimized ecological problems, and production expansion without a capital expenditure for new buildings and equipment. Some precautions are necessary with pre-coated sheet. The product must be handled with more care to prevent scratches and damage to the prefinished surface. Metal finishing of damaged areas is more difficult than on uncoated sheet. Fabrication methods are more restrictive, bend radii must be more generous, and welding practices must be carefully chosen. The basic types of precoating include metallic, pretreated, preprimed, and pre-painted finishing. Metallic coating can be made up to zinc, aluminum, zinc-aluminum alloys, tin, and terne metal. Pretreatment coatings are usually phosphates, and pre-primed finishes can be applied as a variety of organic-type coatings. These can be used as a primed-only coating, or a suitable paint topcoat can be applied. Prepainting consists of applying an organic paint system to steel sheet on a coil coating line either at a mill or at a coil coater. This article will address each of these coating processes. Emphasis will be placed on products that are galvanized by the hot dip process, although much of the discussion is equally applicable to electrogalvanizing and zinc spraying.

Zinc Coatings Galvanizing is a process for rustproofing iron and steel by the application of a metallic zinc coating. It is applicable to products of nearly all shapes and sizes, ranging from nails, nuts, and bolts to large structural assemblies and steel sheet in coils and cut lengths. Other applications include roofing and siding sheets for buildings, silos, grain bins, heat exchangers, hot water tanks, pipe, culverts, conduits, air conditioner housings, outdoor furniture, and mail boxes. On all steel parts, galvanizing provides long-lasting, economical protection against a wide variety of corrosive elements in the air, water, or soil. In the United States, more than 9 × 106 Mg (1 × 107 tons) of steel is produced annually by precoating. A large amount of this total is used by the automotive industry for both unexposed and exposed panels⎯from frames and floor pans to doors, fenders, and hoods (Fig. 1 ). Typically, 75% of the body, chassis, and power train components of one American automobile manufacturer's 1986 models consisted of galvanized precoated sheet (Fig. 2 ). Table 1 indicates that a typical 1986 American car utilized nearly 160 kg (350 lb) of zinc-coated steel components in its material composition. As indicated in Table 2 , undervehicle test coupons evaluated after 2 years of exposure attest to the benefits of precoated steels in combating corrosion (additional information is available in the article "Corrosion in the Automobile Industry" in Corrosion, Volume 13 of ASM Handbook, formerly 9th Edition Metals Handbook. Table 1 Use of zinc-coated steel for a typical 1986 model U.S. car Amount of steel Type One-side galvanized Two-side galvanized

Amount of zinc

kg

lb

kg

lb

33.5

74

0.55

1.21

93

205

3.05

6.72

Zincrometal

29.5

65

0.19

0.41

Net total Source: Ref 2

156

344

3.8

8.34

Table 2 Corrosion of unpainted coated steel test coupons after 2 years of undervehicle exposure

Coating weight per side


Steel thickness

Surface area showing base metal attack, %

Average pit depth Vehicle 1

Vehicle 2

Page 355



01 Sep 2005

µm

mils

µm

mils

14

0

0

15

0.6

27.3

11

0.43

56

2.2

g/m2

oz/ft2

mm

in.

Galvanized 1

120−150

0.39−0.49

0.71

0.028

0.6

Galvanized 2

100−120

0.33−0.39

0.90

0.035

0.3

Galvanized 3

55−90

0.18−0.30

0.45

0.018

0.5

5.0

0

0

15

0.6

0.055

0

1.0

0

0

22

0.87

Material

Vehicle 1 Vehicle 2

Hot dip

Galvannealed 1

80−120

0.26−0.39

1.42

Galvannealed 2

75−85

0.25−0.28

0.89

0.035

0.3

32.8

11

0.43

86

3.4

66

0.22

0.66

0.026

25

56.5

48

1.9

67

2.6

Zn

90

0.30

0.88

0.035

61

86

64

2.5

120

4.7

Zn-15Ni-0.4Co

37

0.12

0.70

0.0275

46

67.5

75

3

81

3.2

Zn-16Ni

20

0.065

0.68

0.027

85

93.5

83

3.3

100

4

Zn-16Ni

40

0.13

0.68

0.027

38

79.3

73

2.9

128

5

Zn-16Al

25

0.08

0.68

0.027

59

84.3

64

2.5

97

3.8

Zn-22Al

40

0.13

0.68

0.027

54

76.5

64

2.5

90

3.5

40

0.13

0.92

0.036

10.8

17.3

53

2.1

73

2.9

>10

>250

>10

One-side galvannealed One-side electrodeposited

Zinc-rich primer One-side Zincrometal Uncoated Cold-rolled steel ... ... 0.51 0.020 100 100 >250 Vehicle 1,660 days, 51,000 km (31,700 miles); vehicle 2,660 days, 53,500 km (33,250 miles). Source: Ref 3

Fig. 1 Use of zinc-coated steels in a 1987 model by one U.S. automaker. Source: Ref 1


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01 Sep 2005

Fig. 2 Pie chart illustrating typical usage of zinc-coated steel components for body, chassis, and power train applications in a 1986 car manufactured by a U.S. automaker. Source: Ref 1


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01 Sep 2005

Metallic zinc is applied to iron and steel by three processes: hot dip galvanizing, electrogalvanizing, and zinc spraying. Most galvanized steel sheet is coated by the hot dip process, although there has been strong growth in electrogalvanizing capacity during the past few years. Corrosion Resistance. The use of zinc is unique among methods for the corrosion protection of steel. The zinc coating serves a twofold purpose: • It protects the steel from corrosive attack in most atmospheres, acting as a continuous barrier shield between the steel and the atmosphere • It acts as a galvanic protector, slowly sacrificing itself in the presence of corrosive elements by continuing to protect the steel even when moderate-sized areas of bare metal have been exposed This latter ability is possible because zinc is more electrochemically active than steel. This dual nature of zinc coatings is also available with some zinc/aluminum alloy coatings, but zinc coatings clearly offer the most galvanic protection. With most protective coatings that act only as a barrier, rapid attack commences when exposure of the base metal occurs. The distance over which the galvanic protection of zinc is effective depends on the environment. When completely and continuously wetted, especially by a strong electrolyte (for example, seawater), relatively large areas of exposed steel will be protected as long as any zinc remains. In air, where the electrolyte is only superficially or discontinuously present (such as from dew or rain), smaller areas of bare steel are protected. The order of magnitude of this throwing power is nominally about 3.2 mm (1=8in.), although this can vary significantly with the type of atmosphere. Nevertheless, galvanized parts exposed outdoors have remained rust free for many years, and the two basic reasons are the sacrificial protection provided by the zinc and the relatively stable zinc carbonate film that forms on the zinc surface to reduce the overall corrosion rate of the zinc coating. The service life of zinc-coated steel is dependent on the conditions of exposure and on the coating thickness, as illustrated in Fig. 3 . Although the coating process used to apply the zinc coating generally does not affect the service life, experience has shown that the corrosion resistance of galvanized coatings in the field cannot be accurately predicted from accelerated laboratory tests. Environmental factors such as atmospheric contaminants (sulfates, chlorides, and so on) and time of wetness have a large influence on the service life of galvanized steel. In polluted areas, such as severe industrial areas, the normally protective zinc carbonate film that forms on the surface of the zinc coating tends to be converted to soluble sulfates that are washed away by rain, thus exposing the zinc to further attack and accelerating the corrosion. Fig. 3 Service life of zinc-coated steel sheet. Service life is measured in years to the appearance of first significant rusting.


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Coating Test and Designations. The thickness (or weight), adhesion, and ductility of a zinc coating can have important effects on its service life and effectiveness against corrosion. Practical tests for these characteristics are described in relevant specifications issued by the American Society for Testing and Materials (ASTM). Tests for coating thickness include microscopic measurement of the cross section, stripping tests in which the coating is removed from a given area (ASTM A 90), electrochemical stripping from a given area (ASTM B 504), and magnetic and electromagnetic methods of measurement (ASTM E 376, A 123, A 754, B 499, and D 1186). Adhesion can be tested and rated by bend test methods described in ASTM A 525 and A 879. Other adhesion test methods include reverse impact and draw bend test. Because the service life of a zinc-coated part in a given atmosphere is directly proportional to the thickness of zinc in the coating (Fig. 3 ), measurement of that amount is very important. The amount of coating is most often measured in terms of weight rather than thickness, usually by the method described in ASTM A 90. Specimens are cut from one or three spots in samples of the sheet, as described in ASTM A 525. These are weighed, the zinc is stripped (dissolved) in an acid solution, and the specimens are reweighed. The weight loss is reported in ounces per square foot of sheet or grams per square meter. When specimens from three spots are checked (triple-spot test), the value of weight loss is the average of the three specimens. When the weight-loss method is used, the amount of coating measured is the total amount on both sides of the sheet. Ordinarily, the zinc coating is applied to both sides of the sheet. Therefore, a 2 oz/ft2coating has 305 g/m2 (1 oz/ft2) on each surface. This 28 g (1 oz) is equivalent to an average thickness of 43 µm (1.7 mils). When zinc-coated sheet is ordered, the minimum amount of coating can be specified as the weight determined by the triple-spot or single-spot test or by coating designations corresponding to these weights (Table 3 ). Table 3 Designations and weights of zinc coating on hot dip galvanized sheet Minimum coating weights(a) Triple-spot test

Single-spot test

Coating designation(b)

g/m2

oz/ft2

g/m2

oz/ft2

G 235

717

2.35

610

2.00

G 210

640

2.10

549

1.80

G 185

564

1.85

488

1.60

G 165

503

1.65

427

1.40

G 140

427

1.40

366

1.20

G 115

351

1.15

305

1.00

G 90

275

0.90

244

0.80

G 60

183

0.60

152

G 01


No minimum

0.50 No minimum

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01 Sep 2005

A 60

183

0.60

152

0.50

A 40

122

0.40

91

0.30

A 01 No minimum No minimum (a) Total weight on both sides of sheet per unit area of sheet. (b) G, regular type of coating; A, zinc-iron alloy type of coating. Source: ASTM A 525

Chromate Passivation. Several types of finishes can be applied to zinc-coated surfaces to provide additional corrosion resistance. The simplest type of finish applicable to fresh zinc surfaces is a chromate passivation treatment. This process is equally suitable for use on hot dip galvanized, electrogalvanized, zinc-sprayed, and zinc-plated articles. Usually, the treatment consists of simply cleaning and dipping the articles in a chromic acid or sodium dichromate solution at about 20 to 30 °C (68 to 86 °F), followed by rinsing in cold fresh water and drying in warm air. The adherent chromate film may have a greenish or greenish-yellow iridescent appearance. Specification ASTM B 201 gives details of tests for measuring the adequacy and effectiveness of the chromate film. Chromate passivation helps prevent staining when galvanized sheet is stored under wet or humid conditions. Therefore, a thin, almost clear chromate or chromate/phosphate passivation film is often applied to the coating on hot dip coating lines. Painting. The selection of galvanized steel as a material for barns, buildings, roofs, sidings, appliances, and many hardware items is based on the sacrificial protection and the barrier coating afforded the base metal by zinc coating. For additional protection and cosmetic appearance, paint coatings are often applied to the galvanized steel. The performance of the coatings is an important economic factor in the durability of this material (Table 4 ). Table 4 Synergistic protective effect of galvanized steel/paint systems in atmospheric exposure Galvanized steel Thickness Type of atmosphere Heavy industrial

Metropolitan (urban)

Marine

µm

mils

Service life(a), years

Paint Thickness µm

mils

Galvanized plus paint Service life(a), years

Thickness µm

mils

Service life(a), years

50

2

10

100

4

3

150

6

19

75

3

14

150

6

5

225

9

29

100

4

19

100

4

3

200

8

33

100

4

19

150

6

5

250

10

36

50

2

19

100

4

4

150

6

34

75

3

29

150

6

6

225

9

52

100

4

39

100

4

4

200

8

64

100

4

39

150

6

6

250

10

67

50

2

20

100

4

4

150

6

36

100

4

40

100

4

4

200

8

66

100 4 40 150 (a) Service life is defined as time to about 5% red rust. Source: Ref 4

6

6

250

10

69

Galvanized steel, both new and weathered, can be painted with a minimum of preparation and with excellent adherence. On hot dip galvanized or zinc-plated steel, it is necessary to use special corrosion-inhibitive primers to prepare the surface before the paint is applied. This is partly because these types of zinc coating are too smooth to provide a mechanical key for the paint or lacquer and partly because the paint appears to react with the unprepared zinc surface in the presence of moisture to weaken the initially formed bond. Many exposure tests have shown that zinc dust-zinc oxide paints (finely powdered zinc metallic and zinc oxide pigment in an oil or alkyd base) adhere best to galvanized steel surfaces under most conditions. Zinc dust-zinc oxide primers can be used over new or weathered galvanized steel and can be top coated with most oil or latex house paints or alkyd enamels. For the maintenance painting of galvanized steel, one or two coats of a zinc dust-zinc oxide paint are often used alone. The paint can be applied by brushing, rolling, or spraying. Zinc sheets to be painted should not be treated at the mill with a chromate treatment, although, they may be given a phosphate treatment to improve the adherence of the paint. Zinc-coated sheet steel is often prepainted in coil form by coil coating, is described in the section "Prepainted Sheet" in this article. Packaging and Storage. Galvanized products in bundles, coils, or stacks of sheets must be protected from moisture, including condensation moisture, until openly exposed to the weather. They must be properly packaged and stored. Otherwise, wet-storage stain, a bulky white deposit that often forms on zinc surfaces stored under wet or humid conditions, may develop. It is important to examine packages of galvanized products for damage and to take prompt action where cuts, tears, or other damage is evident. If the packaging is damaged or if moisture is present, the product should be dried at once and not repiled until thoroughly dry. Erection of materials should begin as soon as possible after the package arrives at the installation site. If temporary storage of the galvanized product is absolutely necessary, it should be indoors. Where indoor storage is not possible, intact waterproof bundles can be stored at the site. The package should be slanted so that any condensation will drain out, and it should be stored sufficiently high to allow air circulation beneath and to prevent rising water from entering. The stacks


Page 360



01 Sep 2005

should be thoroughly covered with a waterproof canvas tarpaulin for protection from rain, snow, or condensation. The use of airtight plastic coverings should be avoided. To deter the formation of wet-storage stain, zinc-coated sheet can be purchased with a mill-applied chromate or chromate/phospate film. Various proprietary mixtures are available. Hot dip galvanizing is a process in which an adherent, protective coating of zinc and iron-zinc alloys is developed on the surfaces of iron and steel products by immersing them in a bath of molten zinc. Most zinc coated steel is processed by hot dip galvanizing. One method of hot dip galvanizing is the batch process, which is used for fabricated steel items such as structurals or pipe. This method involves cleaning the steel articles, applying a flux to the surfaces, and immersing them in a molten bath of zinc for varying time periods to develop a thick alloyed zinc coating. The most common form of hot dip galvanizing for steel sheet is done on a continuous galvanizing line. Coiled sheet is fed from pay-off reels through flatteners. It is then cleaned, bright annealed, and passed through the coating bath. After leaving the coating bath, the coating thickness is controlled by an air knife or steel rolls. The sheet is then cooled and recoiled or cut into lengths. The hot dip process normally coats both sides of the sheet. However, hot dip galvanized sheets can be coated on one side only for special uses, such as automotive exposed panels, by the use of special coating techniques. One-side coated sheet produced by the hot dip process is not commonly available. Continuous coating lines have to be specially modified to make one-side coated product. A typical hot dip galvanized coating produced by the batch process consists of a series of layers (Fig. 4 ). Starting from the base steel at the bottom of the coating, each successive layer contains a higher proportion of zinc until the outer layer, which is relatively pure zinc, is reached. There is, therefore, no real line of demarcation between the iron and zinc, but a gradual transition through the series of iron-zinc alloys that provide a powerful bond between the base metal and the coating. These layers are identified in Table 5 . The structure of the coating (the number and extent of the alloy layers) and its thickness depend on the composition and physical condition of the steel being treated as well as on a number of factors within the control of the galvanizer. Table 5 Properties of alloy layers of hot dip galvanized steels Melting point Layer

Alloy

Crystal structure

Diamond pyramid microhardness

Alloy characteristics

Iron, %

°C

°F

0.03

419

787

Hexagonal

70−72

Soft, ductile

Eta (η)

Zinc

Zeta (ζ)

FeZn13

5.7−6.3

530

986

Monoclinic

175−185

Hard, brittle

Delta (δ)

FeZn7

7.0−11.0

530−670

986−1238

Hexagonal

240−300

Ductile

Gamma (Γ)

Fe3Zn10

20.0−27.0

670−780

1238−1436

Cubic

...

...

1510

2750

Cubic

150−175

Steel base metal Iron

Thin, hard, brittle ...

Fig. 4 Photomicrograph of a typical hot dip galvanized coating. The molten zinc is interlocked into the steel by the alloy reaction, which forms zinc-iron layers and creates a metallurgical bond. 250×

The ratio of the total thickness of the alloy layers to that of the outer zinc coating is affected by varying the time of immersion and the speed of withdrawal of the work from the molten zinc bath. The rate of cooling of the steel after withdrawal is another factor to be considered; rapid cooling gives small spangle size. Sheet galvanizers operating continuous strip processes usually suppress the formation of alloy layers by adding 0.1 to 0.2% Al to the bath; this increases the ductility of the coating, thus rendering the sheet more amenable to fabrication (Fig. 5 ). Other elements can be added to galvanizing baths to improve the characteristics and appearance of the coating. Lead and antimony give rise to well-defined spangle effects. Fig. 5 Microstructure of continuously galvanized steel. In continuous hot dip galvanizing, the formation of various iron-zinc alloy layers is suppressed by the addition of 0.1 to 0.2% Al.


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01 Sep 2005

During batch galvanizing, the zinc-iron alloy portion of the coating will represent 50 to 60% of the total coating thickness. However, certain combinations of elements may result in a coating that is either completely or almost completely alloyed. Visually, the zinc-iron alloy coating will have a gray, matte appearance because of the absence of the free-zinc layer. The free-zinc layer imparts the typical bright finish to a galvanized coating. Because of the greater percentage of the zinc-iron alloy present in the coating, the alloy-type coating may have a lower adherence than the regular galvanized coating. The corrosion resistance of the zinc-iron and free zinc coating types is equal for all practical purposes. Steels containing carbon below 0.25%, phosphorus below 0.05%, and manganese below 1.35% (either individually or in combination) will normally develop regular galvanized coatings when conventional galvanizing techniques are used and when silicon is 0.05% or less or ranges between 0.15 and 0.3%. Fabricators and consumers should be aware that a gray matte appearance may occur in batch galvanizing if silicon content exceeds 0.06%. This matte appearance does not reduce the long-term atmospheric corrosion protection of the galvanized coating. Galvanized coatings on sheet products that are intended to be painted are frequently given treatments to make the spangle less obvious so that it does not show through the paint. A flat spangle without relief (suppressed spangle) can be obtained by small additions of antimony to the molten bath; smaller grain size (minimized spangle) can be produced by spraying the molten zinc with zinc dust, steam, air, or water just before it freezes. Finer grains are less visible through the paint and have narrower and smaller fractures on bending, often permitting the paint to bridge the gap and provide increased protection. Galvanized steel sheet can be temper rolled to flatten surface irregularities such as dross and grain boundaries, thus providing an extra smooth surface more suitable for painting where critical surface requirements exist. At the galvanizing mill, galvanized steel sheet can be given a thermal treatment after coating, which converts all the free zinc to zinc-iron alloy, thus providing a spangle-free surface that is more suitable for painting. It can be painted without pretreatment (but not with all paints). As an added benefit, there is no spangle to show through the paint. However, the zinc-iron alloy coating is somewhat brittle and tends to powder if severely bent in fabrication. Table 6 lists the seven ASTM specifications that cover hot dip galvanized steel sheet products. The general requirements for the products covered in these specifications are described in ASTM A 525. Included in this specification are the bend test requirements given in Table 7 , but not included in these bend test requirements are those for structural (physical) quality sheet, which are given in Table 8 . Table 6 ASTM specifications for hot dip galvanized steel sheet General requirements are given in A 525. Specification

Application or quality

A 361

Sheet for roofing and siding

A 444

Sheet for culverts and underdrains

A 446

Structural (physical) quality sheet

A 526

Commercial quality sheet

A 527

Lock-forming quality sheet

A 528

Drawing quality sheet

A 642

Drawing quality, special killed sheet

Table 7 Bend test requirements for hot dip galvanized steel sheet Table does not apply to structural (physical) quality sheet; see Table 8 instead. Bend diameter for sheet thickness range(a)


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01 Sep 2005

0.33−0.97 mm (0.0131−0.0381 in.)

0.97−1.90 mm (0.0382−0.0747 in.)

1.90−4.46 mm (0.0748−0.1756 in.)

G 235

2t

3t

3t

G 210

2t

2t

2t

G 185

2t

2t

2t

G 165

2t

2t

2t

G 140

t

t

2t

G 115

0

0

t

G 90

0

0

t

G 60

0

0

0

Coating designation

G 01 0 0 0 (a) Value listed is the minimum diameter of rod (or mandrel), in multiples of the galvanized sheet thickness (t) around which the galvanized sheet can be bent 180° in any direction at room temperature without flaking of the coating on the outside of the bend. Source: ASTM A 525

Table 8 Bend test requirements for coating and base metal of structural (physical) quality hot dip galvanized steel sheet Bend diameter for sheet grade(a)

Coating designations or base metal

A

B

G 235

3t

3t

3t

G 210

2t

2t

21=2t

G 185

2t

2t

21=2t

G 165

2t

2t

21=2t

G 140

2t

2t

21=2t

G 115

1

1 =2t

2t

21=2t

G 90

11=2t

2t

21=2t

G 60

11=2t

2t

21=2t

2t

21=2t

1

1 =2t

G 01

C

Base metal 11=2t 2t 21=2t (a) Value listed is the minimum diameter of rod (or mandrel), in multiples of the galvanized sheet thickness (t), around which the sheet can be 180° in any direction at room temperature without flaking of the coating, or cracking of the base metal, on the outside of the bend. There are no bend test requirements for coatings and base metal of grades D, E, and F. Source: ASTM A 446

The typical mechanical properties of commercial quality (CQ), drawing quality (DQ), and drawing quality, special killed (DQSK) hot dip galvanized steel sheet are listed in Table 9 . Commercial quality sheet is satisfactory for applications requiring bending and moderate drawing. Drawing quality sheet has better ductility and uniformity than commercial quality and is excellent for ordinary drawing applications. Drawing quality, special killed sheet is superior to drawing quality and is excellent for applications requiring severe drawing. When higher strength is required, structural quality (SQ) sheet, also called physical quality (PQ) sheet, can be specified, although at some sacrifice in ductility (compare Tables 7 and 8 ). The minimum mechanical properties of structural quality sheet are presented in Table 10 . Additional information is available in the article "Hot Dip Coatings in Corrosion, Volume 13 of ASM Handbook, formerly 9th Edition Metals Handbook. Table 9 Typical mechanical properties of hot dip galvanized, long terne, or aluminized steel sheet Not temper rolled. See Table 10 for properties of structural (physical) quality galvanized sheet. Tensile strength(a) Yield strength(a) MPa

ksi

MPa

ksi

Elongation in 50 mm or 2 in.(a), %

Hardness(a), HRB

310−385

45−56

235−290

34−42

30−38

47−68

340

49

255

37

35

58

DQ

305−350

44−51

220−270

32−39

32−40

42−54

325

47

250

36

37

50

DQ (postannealed)

310−340

45−49

215−260

31−38

37−42

40−52

320

46

235

34

41

46

310−375

45−54

205−270

30−39

34−42

46−58

330

48

240

35

38

52

310−345

45−50

180−230

26−33

38−45

42−55

Quality CQ

DQSK DQSK (postannealed)


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325 (a) Single values below ranges are average values.

47

215

01 Sep 2005

31

40

46

Table 10 Minimum mechanical properties of structural quality hot dip galvanized steel sheet Tensile strength

Yield strength

Grade

MPa

ksi

MPa

ksi

Elongation in 50 mm or 2 in., %

A

310

45

230

33

20

B

360

52

255

37

18

C

380

55

275

40

16

D

450

65

345

50

12

E(a)

570

82

550

F 480 70 345 (a) When the hardness value is 85 HRB or higher, no tensile test is required. Source: ASTM A 446

80

...

50

12

Electrogalvanizing. Very thin formable zinc coatings ideally suited for deep drawing or painting can be obtained on steel products by electrogalvanizing. Zinc is electrodeposited on a variety of mill products: sheet, wire, and, in some cases, pipe. Electrogalvanizing the sheet and wire in coil form produces a thin, uniform coating of pure zinc with excellent adherence. The coating is smooth, readily prepared for painting by phosphatizing, and free of the characteristics spangles of hot dip zinc coatings. Electrogalvanizing can be used where a fine surface finish is needed. The appearance of the coating can be varied by additives and special treatments in the plating bath. Electrodeposited zinc coatings are simpler in structure than hot dip galvanized coatings. They are composed of pure zinc, have a homogeneous structure, and are highly adherent. These coatings are not generally as thick as those produced by hot dip galvanizing. Electrogalvanized coating weights as high as 100 g/m2 (0.3 oz/ft2) have been applied to one or both sides of steel sheet. The normal ranges of coating weights available are listed in ASTM Specifications A 591 and A 879. The coating thicknesses listed in A 591 are typically used when the application does not subject the steel sheet to very corrosive conditions or when the sheet is intended for painting. For more severe corrosion conditions, such as the need to protect cars from road salts and entrapped moisture, heavier coatings in the ranges listed in A 879 are used. These coating weights are applied to the steel sheets used for most body panels. Electrodeposited zinc is considered to adhere to steel as well as any metallic coating. Because of the excellent adhesion of electrodeposited zinc, electrogalvanized coils of steel sheet and wire have good working properties, and the coating remains intact after severe deformation. Good adhesion depends on very close physical conformity of the coating with the base metal. Therefore, particular care must be taken during initial cleaning. Electrodeposition also affords a means of applying zinc coatings to finished parts that cannot be predipped. It is especially useful where a high processing temperature could damage the part. One advantage of electrodeposition is that it can be done cold and therefore does not change the mechanical properties of the steel. Zincrometal is also used for outer body panels in automobiles. First introduced in 1972, Zincrometal is a coil coated product consisting of a mixed-oxide underlayer containing metallic zinc particles and a zinc-rich organic (epoxy) topcoat. It is weldable, formable, paintable, and compatible with commonly used adhesives. Zincrometal is primarily used in one-side applications to protect against inside-out corrosion. The corrosion resistance of Zincrometal is not as good as that of hot dip galvanized steels (Ref 1), and its use is declining substantially as more electrogalvanized steels and other types of coatings are employed. Zinc alloy coated steels have also been developed. Coatings include zinc-iron (15 to 80% Fe) and zinc-nickel (10 to 14% Ni) alloys. These coatings are applied by electrodeposition. Zinc-iron coatings offer excellent corrosion resistance and weldability. Zinc-nickel coatings are more corrosion resistant than pure zinc coatings, but problems include brittleness from residual stresses and the fact that the coating is not completely sacrificial, as is a pure zinc coating. This can led to accelerated corrosion of the steel substrate if the coating is damaged (Ref 5). Multilayer coatings that take advantage of the properties of each layer have been developed in Europe. An example of this is Zincrox, a zinc-chromium-chromium oxide coating (Ref 5). The CrOx top layer of this coating acts as a barrier to perforation and provides excellent paint adhesion and weldability (Ref 5). Another relatively new development in zinc alloy coatings is Galfan, a Zn-5Al-mischmetal alloy coating applied by hot dipping. Applications in the United States are limited, but European automakers have used Galfan in such applications as brake servo housings, headlight reflectors and frames, and universal joint shrouds (Ref 6). Galfan is also being considered for oil pans, and heavily formed unexposed body panels. Detailed information is available in the article "Electroplated Coatings" in Corrosion, Volume 13 of ASM Handbook, formerly 9th Edition Metals Handbook. Zinc spraying consists of projecting atomized particles of molten zinc onto a prepared surface. Three types of spraying pistols are in commercial use: the molten metal pistol, the powder pistol, and the wire pistol. The sprayed coating is slightly rough and slightly porous; the specific gravity of a typical coating is approximately 6.35, compared to 7.1 for cast zinc. This slight porosity does not affect the protective value of the coating, because zinc is anodic to steel. The zinc corrosion products that form during service fill the pores of the coating, giving a solid appearance. The slight roughness of the surface makes it an ideal base for paint, when properly pretreated. On-site spraying can be performed on finished parts of almost any shape or size. When applied to finished articles, welds, ends and rivets receive adequate coverage. Moreover, it is the only satisfactory method of depositing unusually thick zinc coatings


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(≥0.25 mm, or 0.01 in.)

Aluminum Coatings Aluminized (aluminum-coated) steel sheet is used for applications where heat resistance, heat reflectivity, or barrier-layer resistance to corrosion is required. Aluminum coating of steel sheet is done on continuous lines similar to those used for hot dip galvanizing of steel sheet. Cold-rolled steel sheet is hot dipped into molten aluminum or an aluminum alloy containing 5 to 10% Si. The coating consists of two layers, the exterior layer of either pure aluminum or aluminum-silicon alloy and the steel base metal, with an aluminum-iron-silicon alloy layer in between. The thickness of this alloy can significantly affect the ductility, adhesion, uniformity, smoothness, and appearance of the surface and is controlled for optimum properties. Aluminum-coated sheet steel combines the desirable properties of aluminum and steel. Steel has a greater load-bearing capacity, having a modulus of elasticity of about three times that of unalloyed aluminum. The thermal expansion of steel is approximately half as much as that of aluminum. The aluminum coating offers corrosion resistance, resistance to heat and oxidation, and thermal reflectivity. Typical applications are: • • • • • • • • • • • • • • •

Automotive mufflers and related components Catalytic converter heat shields Drying and baking ovens Industrial heating equipment Fireplaces Home incinerators and furnaces Fire and garage doors Kitchen and laundry appliances Metal buildings Agricultural equipment Silo roofs Playground equipment Outdoor furniture Signs, masts, and lighting fixtures Containers and wrappers

Coating Weight. Aluminum coatings on steel sheet are designated according to total coating weight on both surfaces in ounces per square foot of sheet, as indicated in Table 11 . These coating categories are listed in ASTM Specification A 463. Type 1, Light Coating, is recommended for drawing applications and when welding is a significant portion of the fabrication. Type 1, Regular or Commercial, has approximately a 25 µm (1 mil) thick coating on each surface (Fig. 6 a). It is designated for applications requiring excellent heat resistance. Type 2 has a coating approximately 50 µm (2 mil) thick on each side (Fig. 6 b). It is frequently used for atmospheric corrosion resistance. Coating weight on specimens from aluminum-coated sheet is determined by the test method in ASTM A 428. Figure 7 demonstrates how a typical rear suspension of a front-wheel drive vehicle utilizes type 1 aluminized steel components having a coating of Al-9Si-3Fe in conjunction with galvanized front pivot hangers, mounting brackets, and braces. Table 11 Designations and weights of aluminum coating on aluminized steel sheet Minimum coating weight(a) Triple-spot test

Single-spot test

Coating designation

g/m2

oz/ft2

g/m2

oz/ft2

T1 25 (light)

80

0.25

60

0.20

T1 40 (regular)

120

0.40

T2 (regular) 230 0.75 (a) Total weight on both sides of sheet per unit area of sheet. Source: ASTM A 463 and A 428

90

0.30

200

0.65

Fig. 6 Microstructure of aluminum coatings on steel. Left: Type 1 coating from top: a nickel filler, aluminum-silicon alloy, aluminum-silicon-iron alloy, and steel base metal. Right: Type 2 coating forms a layer of essentially pure aluminum (top) with scattered gray particles of aluminum-iron; the light gray center layer is aluminum-iron, and the bottom layer is the base steel. Both 1000×


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Fig. 7 Typical galvanized and aluminized steel rear suspension components used in American front-wheel drive automobiles. Source: Ref 1

Base Metal and Formability. Aluminum coating can be applied to CQ, DQ, or DQSK steel sheet, depending on the severity of the forming or drawing required. Only moderate forming and drawing are recommended for aluminized steel sheet, but there are numerous intricate components for heating, combustion, and other equipment being produced. Shallow crazing (hairline cracks) may occur in the coating if the bending and forming are too severe. To eliminate crazing, the radius of the bend should be increased. If the crazing is deep enough to expose the steel to the atmosphere during service, staining may occur. These stains generally have minimal effect on the serviceability of the product, because the corrosion stops at the crazed area after a relatively short exposure period. However, if water collects and does not drain off, corrosion products are dissolved and corrosion continues. The mechanical properties of hot dip aluminized steel sheet are essentially the same as those of hot dip galvanized steel sheet (Table 9 ). When high strength is required, SQ aluminized steel sheet may be specified, although at some sacrifice in ductility. Corrosion Resistance. The value of aluminum as a protective coating for steel sheet lies principally in its inherent corrosion resistance. In most environments, the long-term corrosion rate of aluminum is only about 15 to 25% that of zinc. Generally, the protective value of an aluminum coating on steel is a function of coating thickness. The coating tends to remain


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intact and therefore provides long-term protection. Aluminum coatings do not provide sacrificial protection in most environments, particularly in atmospheric exposure. This is because a protective oxide film forms on the coating, which tends to passivate the aluminum and retard sacrificial protection. If the oxide film is destroyed, the aluminum will provide sacrificial protection to the base metal. In marine or salt-laden environments, the aluminum coating will protect sacrificially wherever chlorides destroy the surface oxide film. Although staining or light rusting of the steel may occur at cut edges or crazing may occur where the aluminum does not protect, this action diminishes with further exposure time because of the self-sealing action of corrosion products. However, if insufficient slope or drainage permits water to pond or remain instead of running off freely, the corrosion products are dissolved and rusting will continue. Heat Resistance. Aluminum-coated sheet steel has excellent resistance to high-temperature oxidation. At surface temperatures below about 510 °C (950 °F), the aluminum coating protects the steel base metal against oxidation without discoloration. Between 510 and 675 °C (950 and 1250 °F), the coating provides protection to the steel, but some darkening may result from the formation of aluminum-iron-silicon alloy. The alloy is extremely heat resistant, but upon long exposure at temperatures above 675 °C (1250 °F), the coating may become brittle and spall because of a different coefficient of expansion from that of the steel. Because of their good resistance to scaling, combined with the structural strength of the steel base metal, type 1 coatings are used in automotive exhaust systems, heat exchangers, ovens, furnaces, flues, and tubing. The higher strength of the steel base metal, which melts at 1580 °C (2876 °F), enables steel sheet coated with either type 1 or type 2 coatings to perform for a longer time than aluminum alone in the event of fire. Heat Reflection. The thermal reflectivity of aluminum-coated steel sheet is comparable to that of aluminum sheet. It is superior to galvanized steel sheet after a relatively short exposure time. All three sheet materials have thermal reflectivity of approximately 90% before exposure. However, after a few years, the value for galvanized steel decreases more than that for aluminized steel. Aluminum and aluminum-coated steel sheet retain 50 to 60% of their heat reflectivity. This is advantageous where heat must be confined, diverted, or excluded, as in heat transfer applications. When used for roofing and siding, aluminum-coated sheet keeps buildings cooler in summer and warmer in winter. Weldability. Aluminum-coated steel sheet can be joined by electric resistance welding (spot welding or seam welding). It can also be metal arc welded, flash welded, or oxyacetylene welded. Thorough removal of grease, oil, paint, and dirt followed by wire brushing is recommended before joining. Special fluxes are required for metal arc and oxyacetylene welding. During spot welding, electrodes tend to pick up aluminum, and the tips must be dressed more frequently than during spot welding of uncoated steel. Also, higher current density is required. Painting is generally unnecessary, but aluminum-coated sheet steel can be painted similarly to aluminum sheets. This includes removal of oil or grease and treatment with a phosphate, chromate, or proprietary wash-type chemical before painting. Handling and Storage. The coating on aluminized steel sheet is soft, and care should be taken to avoid scratching and abrasion of the soft coating, which will mar the appearance and allow staining if the coating is removed. Wet-storage stains develop on aluminum-coated steel sheet that is continuously exposed to moisture. The sheet should be inspected for entrapped moisture when received and then stored indoors in a warm, dry place. Some added protection can be obtained by ordering the sheet oiled or chemically treated for resistance to wet-storage stain.

Aluminum-Zinc Alloy Coatings In recent years, the desire and need to improve the corrosion resistance of galvanized coatings while retaining sacrificial galvanic corrosion behavior at sheared edges, and so on, have led to the commercial development of two types of hot dip aluminum-zinc alloy coatings. One type consists of about 55% Al and 45% Zn; the other type, zinc plus 5% Al. Both coating types contain small amounts of other alloying elements to improve wettability and/or coating adhesion during forming. Descriptions of these products are contained in ASTM A 792 (55Al-45Zn coating) and A 875 (Zn-5A1 coating). These specifications include the general requirements, the coating categories available, and the product types that are available. The 55% Al coating has been produced worldwide by a number of steel companies for more than 10 years. Its primary use is for preengineered metal buildings, especially roofing. In most environments, this coating has been found to have two to four times the corrosion resistance of galvanized coatings while retaining an adequate level of galvanic protection to minimize the tendency toward rust staining at edges and other breaks in the coating. Figure 8 illustrates the corrosion resistance of 55Al-Zn-coated steel exposed to four atmospheres. The coated sheet is available in similar grades (CQ, DQ, high strength, and so on) as hot dip galvanized and can be subjected to similar types of forming. It can also be painted either by coil-line painting methods or postpainting after fabrication. Fig. 8 Coating thickness loss of 55Al-Zn-coated steel in four atmospheres. Source: Ref 7


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The coating microstructure consists of an aluminum-iron intermetallic alloy bond between the steel and outer coating-metal layer (Fig. 9 ). This outer coating layer has a duplex microstructure, a matrix phase of an aluminum-rich composition, and a zinc-rich interdendritic phase. This zinc-rich phase corrodes preferentially to provide the galvanic corrosion protection. The coating contains about 2% Si, which is present in the microstructure as an elemental silicon phase. The silicon is added only to inhibit growth of the alloy layer during the hot dip coating operation. Fig. 9 Microstructure of an aluminum-zinc coated sheet. Etched with Amyl nital. 500×

Although this 55% Al coating is primarily used for metal-building applications, there are a variety of other applications, including appliances and automotive parts. It offers a level of heat-oxidation resistance intermediate between galvanized and aluminized coatings. The Zn-5Al coating is also produced worldwide, but it is not as commonly available as the 55% Al coating. Its primary attribute is improved coating ductility compared to hot dip galvanized coatings. This feature, along with a somewhat improved corrosion resistance, makes this coated-sheet product attractive for deep-drawn parts. Also, for prepainted sheets such as roll-formed metal-building panels, the improved coating ductility minimizes the tendency toward cracking of the paint along tension bends. The Zn-5Al coated sheet is also available in similar grades (CQ, DQ, and so on) as hot dip galvanized. It is readily paintable, including coil-line prepainting. Both types of aluminum-zinc coating have features that make them more attractive than galvanized for certain applications. Selection of either one should be based on consideration of the desired attributes and differences in fabricability, weldability, paintability, and so on, compared to the other coatings available.


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Tin Coatings Tin coatings are applied to steel sheet by electrolytic deposition or by immersion in a molten bath of tin (hot dip process). Hot dip tin coatings provide a nontoxic, protective, and decorative coating for food-handling, packaging, and dairy equipment, and they facilitate the soldering of components used in electronic and electrical equipment. In the United States, hot dip tin coating has been replaced by electrolytic tin coating. Electrolytic tin coated steel sheet is used where solderability, appearance, or corrosion resistance under certain conditions is important, as in electronic equipment, food-handling and processing equipment, and laboratory clamps. It is generally produced with a matte finish formed by applying the coating to base metal sheet called black plate, which has a dull surface texture, and by leaving the coating unmelted. It can also be produced with a bright finish by applying the coating to base metal having a smooth surface texture and then melting the coating. Electrolytic tin coated sheet is usually produced in nominal thicknesses from 0.38 to 0.84 mm (0.015 to 0.033 in.) and in widths from 305 to 915 mm (12 to 36 in.). Electrolytic tin coated steel sheet can be specified to one of the five coating-weight designations listed in Table 12 . The coating weight is the total amount of tin on both surfaces, expressed in ounces per square foot of sheet area. Electrolytic coatings can be applied to CQ, DQ, or DQSK steel sheet, depending on the severity of the forming or drawing required. They can also be applied to SQ steel sheet when higher strength is required. Electrolytic tin coated steel sheet is covered in ASTM A 599. The mechanical properties of the steel sheet are unchanged by the electrolytic tin coating process. Table 12 Designations and weights of tin coating on electrolytic tin coated steel sheet Minimum coating weight(a) Coating designation

Triple-spot test

Single-spot test

g/m2

oz/ft2

g/m2

oz/ft2

3.7

0.012

2.8

0.009

50

7.3

0.024

5.6

0.018

75

11.0

0.036

8.2

0.027

100

14.6

0.048

11.0

0.036

125 18.3 0.060 (a) Total weight on both sides of sheet per unit area of sheet. Source: ASTM A 599

13.8

0.045

25

Terne Coatings Long terne steel sheet is carbon steel sheet continuously coated by the hot dip process with terne metal (lead with 3 to 15 wt% Sn). This coated sheet is duller in appearance than tin-coated sheet, hence the name (terne) from the French, which means dull or tarnished. The smooth, dull coating gives the sheet corrosion resistance, formability, excellent solderability, and paintability. The term long terne is used to describe terne-coated sheet, while short terne is used for terne-coated plate. Short terne, also called terneplate, is no longer produced in the United States. Because of its unusual properties, long terne sheet has been adapted to a wide variety of applications. Its greatest use is in automotive gasoline tanks. Its excellent lubricity during forming, solderability and special corrosion resistance make the produce well suited for this application. Other typical applications include: • Automotive parts, such as air conditioners, air filters, cylinder head covers, distributor tubes, oil filters, oil pans, radiator parts, and valve rocker arm covers • Electronic chassis and parts for radios, tape recorders, and television sets • File drawer tracks • Fire doors and frames • Furnace and heating equipment parts • Railroad switch lamps • Small fuel tanks for lawn mowers, power saws, tractors, and outboard motors Long terne sheet is often produced to ASTM A 308. The coatings are designated according to total coating weight on both surfaces in ounces per square foot of sheet area, as indicated in Table 13 . For applications requiring good formability, the coating is applied over CQ, DQ, or DQSK low-carbon steel sheet. The terne coating acts as a lubricant and facilities forming, and the strong bond of the terne metal allows it to be formed along with the base metal. When higher strength is required, the coating can be applied over SQ low-carbon steel sheet, although there is some sacrifice in ductility. In general, the mechanical properties of hot dip terne-coated steel are similar to those for cold-rolled steel. Terne coatings are applied by a flux-coating process that does not include in-line annealing. Therefore, the mechanical properties are obtained by pre-annealing using cycles comparable to those used for cold-rolled sheet.


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Table 13 Designations and weights of lead-tin coating on terne-coated steel sheet Minimum coating weight(a) Coating designation

Triple-spot test g/m2

LT01

Single-spot test oz/ft2

g/m2

No minimum

oz/ft2 No minimum

LT25

76

0.25

61

0.20

LT35

107

0.35

76

0.25

LT40

122

0.40

92

0.30

LT55

168

0.55

122

0.40

LT85

259

0.85

214

0.70

LT110 336 1.10 (a) Total weight on both sides of sheet per unit area of sheet. Source ASTM A 308

275

0.90

Lead is well known for its excellent corrosion resistance, and terne metal is principally lead, with some tin added to form a tight, intermetallic bond with steel. The excellent corrosion resistance of terne sheet accounts for its wide acceptance as a material for gasoline tanks. However, because lead does not offer galvanic protection to the steel base metal, care must be exercised to avoid scratches and pores in the coating. Small openings may be sealed by corrosion products of iron, lead, and oxygen, but larger ones can corrode in an environment unfavorable to the steel base metal. Long terne sheet can be readily soldered with noncorrosive fluxes using normal procedures because the sheet is already presoldered. This makes it a good choice for television and radio chassis and gasoline tanks, for which ease of solderability is important. It can also be readily welded by either resistance seam or spot welding methods. However, when the coating is subjected to high temperatures, significant concentrations of lead fumes can be released. Because of the toxicity of lead, the Occupational Safety and Health Administration and similar state agencies have promulgated standards that must be followed when welding, cutting, or brazing metals that contain lead or are coated with lead or lead alloys. Long terne sheet has excellent paint adherence, which allows it to be painted using conventional systems, but this product is not usually painted. When painting is done, no prior special surface treatment or primer is necessary, except for the removal of ordinary dirt, oil, and grease. Oiled sheet, however, should be thoroughly cleaned to remove the oil. Long terne sheet is normally furnished dry and requires no special handling. It should be stored indoors in a warm, dry place. Unprotected, outdoor storage of coils or bundles can result in white or gray staining of the terne coating, and if there are pores in the terne coating, rust staining can occur.

Phosphate Coatings The phosphate coating of iron and steel consists of treatment with a dilute solution of phosphoric acid and other chemicals by which the surface of the metal, reacting chemically with the phosphoric acid, is converted into an integral layer of insoluble crystalline phosphate compound. This layer is less reactive than the metal surface and at the same time is more absorbent of lubricants or paints. Because the coating is an integral part of the surface, it adheres to the base metal tenaciously. The weight and crystalline structure of the coating, as well as the extent of penetration of the coating into the base metal, can be controlled by the method of cleaning before treatment, the method of applying the solution, the duration of treatment, and the changes in the chemical composition of the phosphating solution. The two types of phosphate coatings in general use are zinc phosphate and iron phosphate. Within each type, chemical composition can be modified to suit various applications. When zinc phosphate coatings are mill applied to galvanized sheets, the sheets are ready for immediate painting with the many paints readily available from industrial and retail suppliers. The zinc phosphate coated product is often referred to as phosphatized. Minor cleaning with mineral spirits, paint thinner, or naphtha may be necessary to remove fingerprints, oils, or dirt picked up during fabrication or handling. When mill-phosphatized sheets that are to be baked after painting are exposed to humid storage conditions for long periods of time, prebaking for several minutes at 150 °C (300 °F) prior to painting may be required to prevent blistering during baking. The chief application for iron phosphate coatings is as a paint base for uncoated carbon steel sheet. Such a coating can be applied on coil coating lines. The greatest tonnage of phosphate-coated steel is low-carbon flat-rolled material, which is used for applications such as sheet metal parts for automobiles and household appliances. Applications of the coatings range from simple protection to prepaint treatments for painted products, such as preengineered building panels and the side and top panels of washing machines, refrigerators, and ranges. Phosphate coatings require a clean surface. The cleaning stage preceding phosphating removes foreign matter from the surface and makes uniformity of coating possible. This involves removal of oils, greases, and associated dirt by solvent degreasing or alkaline cleaning followed by thorough rinsing. Phosphate coatings are applied by spray, immersion, or roller coating. A phosphate coating beneath a paint film helps prevent moisture and other corrosives that may penetrate the paint from reaching the metal. This prevents or delays the electrochemical reactions that lead to corrosion or rust. If the paint film sustains scratches or damage that exposes bare metal, the phosphate coating confines corrosion to the exposed metal surface, preventing


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the corrosion from spreading underneath the paint film. In painting applications, coarse or heavy phosphate coatings may be detrimental; they can absorb too much paint, thus reducing both gloss and adhesion, especially if deformation of the painted sheet steel is involved.

Preprimed Sheet Primer paint coats are frequently applied to steel sheet at the mill or by a coil coater. Because their purpose is corrosion protection, they contain corrosion-inhibiting substances such as zinc powder, zinc chromate, or other compounds of zinc and/or chromium. Preprimed sheets are especially useful for parts that will have limited access after fabrication, rendering coating difficult. Parts made from preprimed sheet may receive a top coat after fabrication. The mill-applied phosphate coatings described in the previous section can also be considered prepriming treatments. Formability (Ref 8). Preprimed steel offers advantages in forming metal fabrication through: • • • •

Consistent surface morphology Reduced surface friction (reducing the flow over die surfaces) and reduced die wear, especially on the binder surfaces Reduced flaking and powdering (requiring less die maintenance), reduced need for metal finishing, and fewer surface defects Reduced galling

The painted surface acts as a cushion between substrate and stamping dies, which lessens the need for in-die lubrication and extends the life of the stamping die. The preprimed, prepainted surface can withstand severe forming and stretching. Thus, the need for lubricant is reduced or eliminated. This is turn provides a clean process environment and reduces the need for extensive cleaning along with phosphating and electrocoating. Zinc Chromate Primers. Zinc chromate pigments are useful as corrosion inhibitors in paint. They are used as after-pickling coatings on steel and in primers. Federal specifications TT-P-57 and TT-P-645 cover zinc chromate paints. Zinc chromate pigments are unique; they are useful as corrosion inhibitors for both ferrous and nonferrous metals. Zinc-Rich Primers. In recent years, manufacturers have developed various priming paints that will deposit films consisting mainly of metallic zinc that have many properties in common with the zinc coatings applied by hot dip galvanizing, electroplating, metal spraying, or mechanical plating methods. Such films will provide some degree of sacrificial protection to the underlying steel if they contain 92 to 95% metallic zinc in the dry film and if the film is in electrical contact with the steel surface at a sufficient number of points. The type of zinc dust used in protective coatings is a heavy powder, light blue-gray in color, with spherically shaped particles having an average diameter of approximately 4 µm (160 µin.). Such powder normally contains 95 to 97% free metallic zinc with a total zinc content exceeding 99%. Many zinc-rich paints are air drying, although oven-curable primers containing a high content of zinc dust are available. Depending on the nature of the binder, zinc-rich primers are classified as inorganic or organic. The inorganic solvent-base types are derived from organic alkyl silicates, which, upon curing, become totally inorganic. The organic zinc-rich coatings are formed by using zinc dust as a pigment in an organic binder. This binder may be any of the well-known coating vehicles, such as chlorinated rubber or epoxy. The zinc dust must be in sufficient concentration so that the zinc particles are in particle-to-particle contact throughout the film. In this way, zinc provides cathodic protection to the base metal. With the organic binder, there is no chemical reaction with the underlying surface, but the organic vehicle must wet the surface thoroughly to obtain mechanical adhesion. The inorganic zinc coating forms its film and adheres to the steel surface by quite different means. The chemical activity during coating is quite similar for either water-or solvent-base inorganic binders. Zinc is the principal reactive element in the inorganic systems and is primarily responsible for the development of initial insolubility. Zinc-rich primers offer a more versatile application of zinc to steel than galvanizing. Large, continuous, complex shapes and fabricated new or existing structures can be easily coated at manufacturing shops or in the field. The performance of zinc-rich primers has earned them a prominent place in the field of corrosion protection coatings. For example, zinc-rich primers are being preapplied to steel sheet as the first coat of a two-coat system for appliance applications such as refrigerator liners. However, the limitations of zinc-rich paints include cost and the required cleanliness of steel surfaces. They must be top coated in severe environments (pH under 6.0 and over 10.5). The following comparisons should be helpful in selecting the binder system that is most suitable for an application. Inorganics have superior solvent and fuel resistance. They can withstand temperatures to 370 °C (700 °F) and are much easier to clean up after use. Inorganics do not blister upon exposure and are unaffected by weather, sunlight, or wide variations in temperature. They do not chalk, peel, or lose thickness over long periods of time. Also, they are easier to weld through and have excellent abrasion resistance and surface hardness. Organics use chlorinated rubber, epoxy, vinyl, phenoxy, or other coating vehicles, and the properties of the system are based on the characteristics of the vehicle used.

Organic Composite Coatings Organic composite coated steels have been developed mainly by Japanese steelmakers in cooperation with automakers in that country, although development is underway in other countries as well. These coil coated products generally employ an electroplated zinc alloy base layer and a chemical conversion coating under a thin organic topcoat containing a high percentage


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of metal powder (Ref 9, 10, 11). The thinness of the organic topcoat allows for good formability without the risk of damaging the coating. Figure 10 compares the corrosion resistance of one of these organic composite coated sheet steels to cold-rolled steel and to Zincrometal. Another of these products uses an organic-silicate composite topcoat only about 1 µm (40 µin.) thick and has corrosion resistance and weldability superior to that of Zincrometal (Ref 10). A bake-hardenable version of this material has also been developed (Ref 10). Researchers at a third Japanese steel company have developed a bake-hardenable organic composite coated sheet steel with a 0.8 to 1.5 µm (32 to 60 µin.) thick organic topcoat. The material possesses corrosion resistance, formability, and weldability equivalent to that of Zincrometal-KII, which uses a 7 µm (280 µin.) thick topcoat (Ref 11). Production of these composite-coated materials is increasing in anticipation of increased demand from Japanese automakers. Fig. 10 Corrosion of heavily worked samples of a composite-coated steel, Zincrometal, and cold-rolled steel in a laboratory cyclic test. Test consisted of 28-min cycles of dipping in 5% saline solution at 40 °C (100 °F), humidifying at 50 °C (120 °F), and drying at 60 °C (140 °F). Source: Ref 9

A similar material has been developed in the United States. This material has an electrodeposited zinc alloy base coat, a mixed intermediate layer of chromium oxide and zinc dust, and an organic topcoat for barrier protection (Ref 12). Figure 11 is a micrograph showing the cross section of the composite-coated steel. In salt spray tests comparing this material to electrodeposited zinc-nickel and Zincrometal, zinc-nickel failed after 216 h, Zincrometal at 480 h, and the composite coating at 960 h (Ref 12). This material was developed to have weldability, formability, and adhesive compatibility similar to that of Zincrometal. Developmental work continuing. Fig. 11 Scanning electron micrograph of cross section through a composite-coated sheet steel. Source: Ref 12

Organic-Silicate Composite Coatings (Ref 13), Zinc-nickel electroplated steel sheet coated with an organic-silicate composite was developed by a Japanese steel company in an attempt to combine a highly corrosion resistant base zinc-nickel coating with a protective surface layer to prolong the coating life. With a view to forming a thin film with high corrosion resistance, the protective layer was designed as a two-layer protective film structure composed of a chromate film as a lower layer and the organic-silicate composite coating (the composite resin) as an upper layer. This protective film structure improves the corrosion resistance not only by the individual effects of each layer, such as the passivation of chromate film and the excellent corrosion resistance of the composite resin, but also by the suppression of excessive dissolution of Cr6+ from the lower chromate film layer by the sealing effect of the upper composite resin layer. This sealing effect sustains the passivation of chromate film more effectively in the corroding environment.

Prepainted Sheet


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Prepainted steel sheet is a large and rapidly expanding market. The sheet is coated in coil form in a continuous coil-painting facility. Lower production costs, improved product quality, elimination of production hazards in the shop, customer satisfaction, conservation of energy, elimination of ecological problems, and the ability to expand production without capital expenditure for new buildings and equipment are some of the advantages of prepainted sheet over postpainting. Fabricated parts are readily joined by indirect projection welding, adhesives, tabs, and fasteners. Typical applications of prepainted steel sheet include tool sheds, preengineered buildings, swimming pools, automobiles, lighting fixtures, baseboard heaters, truck vans, mobile homes, home siding, metal awnings, air conditioners, freezer, refrigerators, ranges, washers, and dryers. Selection of Paint System. A wide variety of paint systems are available on prepainted sheet. In selecting the proper system for a particular application, the user must consider fabrication requirements, the service life desired, and the service conditions, that will be encountered. For example, in an aggressive environment a plastisol coating may be required. For a deep draw, a vinyl coating should be used instead of a polyester. For resistance to fading in sunlight, a silicone polyester may be suggested instead of a polyester or a vinyl paint. In the preengineered building field, the paint system must be capable of being roll formed and still perform over the years under a wide variety of conditions without chalking, fading, cracking, or blistering. In the automotive field, the drawing properties of the coating must be considered in addition to corrosion protection from road salts. In the appliance industry, a high-gloss finish that will bend without cracking is important, along with resistance to such materials as detergents, solvents, mustard, ketchup, shoe polish, grape juice, and grease. Other product requirements frequently considered when selecting an appliance paint are color, hardness, adhesion, resistance to abrasion, corrosion, humidity, heat, and pressure marking. For severe corrosion service and decorative effects, heavier coatings are supplied, often by laminating or bonding a solid film to the metal substrate. Typical applications include buildings, roofing and siding near pickling tanks, chemicals and other corrosive environments, and storm drains and culverts that are subjected to corrosive soils, mine acids, sewage, and abrasion. These culvert coatings can be a thermoplastic coal tar-base laminate 0.3 to 0.5 mm (0.012 to 0.020 in.) thick, or they can be a film of polyvinyl chloride. Design Considerations. In using prepainted sheet, design should be considered. If necessary, binding radii, location of exposed edges, fastening methods, welding techniques, corner assembly, and other features should be modified to make them compatible with the base metal and paint system. For example, if a polyester paint is applied to bare steel sheet, a minimum bend radius of 3.2 mm (1=8in.) is suggested to minimize cracking and crazing of the paint. If hot dip galvanized sheet is the substrate, of minimum bend radius of 6.4 mm (1=4in.) should be used instead. Otherwise, the zinc coating may crack with sharper bending, and the paint may not be elastic enough to bridge the crack. Paint is often cured at temperatures as high as 240 °C (465 °F). At the higher paint curing temperatures, the steel sheet may become fully aged and cause yield point elongation to return. The sheet is subject to the formation of stretcher strains during subsequent forming. Normally, return of yield point elongation is not objectionable in these applications. However, the formed part will sometimes be given a critical amount of strain, and strain lines may become visible. Frequently, this problem can be overcome by proper shop practices, particularly if the part has been roll formed. At times, however, it is necessary to use killed steels, which are considered essentially nonaging. Shop Practices. Because a prepainted surface is composed of an organic material, the abuse that this surface can withstand is less than that of a metal sheet surface. Therefore, prior to using prepainted sheet for the first time, it is advisable to train shop personnel in proper handling practices and to examine shop equipment to eliminate sources of scratches. For example, dies, brake presses, and roll-forming equipment must have highly polished surfaces free of gouges, score marks, and so on. Clearances of the dies must be such that wiping of the paint film is avoided. Similarly, some care is needed when formed parts are put on carts or in containers for transfer from one location to another. It is not acceptable simply to pile one part on top of another. Good housekeeping is important to minimize the source of scratches. Frequent reexamination of shop equipment and parts containers is necessary to minimize scratches. Handling scratches can be refinished by retouching, which is costly and time consuming. Packaging and Handling. Shop and field conditions should be considered when selecting packaging for prepainted sheet. Transit pickoff and job-site corrosion from entrapped moisture can be serious problems. For preengineered building sheets, for example, packaging after roll forming should include waterproof paper (no plastic wrapping), support sheets to prevent sagging, and pressure boards. Mixing sheets of different lengths in the bundle should be avoided. Once the bundle of formed prepainted sheets arrives at the job site, it should be inspected to determine if the packages are still intact and resistant to the weather. Wherever possible, sheets should be erected on the day of delivery, or they should be protected from water condensation. Under-roof storage is desirable. However, if this is not possible, the waterproof bundles should be slanted so that any condensation will drain out. Damaged packages should be opened, inspected, and the sheets separated to allow complete drying. In addition to the prevention of moisture entrapment described above, chips from drilling operations should be brushed away to prevent rust spotting. Prepainted sheets should be installed with corrosion-resistant fasteners. The installation of sheets that are in contact with the soil should be avoided. Oil, grease, fingerprints, and other contaminants should be removed after installation. REFERENCES 1. D.J. Bologna, Corrosion Resistant Materials and Body Paint Systems for Automotive Applications (SAE Paper 862015), in Proceedings of the Automotive Corrosion and Prevention Conference, P-188, Society of Automotive Engineers, 1986, p 69−80 2. "US Automotive Market for Zinc Coatings 1984−1986," Zinc Institute Inc. 3. R. J. Neville and K.M. DeSouza, Electrogalvanized or Hot Dip Galvanized⎯Results of Five Years of Undervehicle


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01 Sep 2005

Corrosion Testing (SAE Paper 862010), in Proceedings of the Automotive Corrosion and Prevention Conference, P-188, Society of Automotive Engineers, 1986, p 31−40 4. J.F.H. van Eijnsbergen, Supplement (to Twenty Years of Duplex Systems), Thermisch Verzinken, Vol 8, 1979 5. M. Memmi et al., A Qualitative and Quantitative Evaluation of Zn + Cr-CrOx Multilayer Coating Compared to Other Coated Steel Sheets (SAE Paper 862028), in Proceedings of the Automotive Corrosion and Prevention Conference, P-188, Society of Automotive Engineers, 1986, p 175−185 6. R.F. Lynch and F.E. Goodwin, "Galfan Coated Steel for Automotive Applications," SAE Paper 860658, Society of Automotive Engineers, 1986 7. H.E. Townsend and J.C. Zoccola, Atmospheric Corrosion Resistance of 55% Al-Zn Coated Sheet Steel: 13-Year Test Results, Mater. Perform., Vol 18, 1979, p 13−20 8. B.K. Dubey, Prepainted Steel for Automotive Application, in Corrosion-Resistant Automotive Sheet Steels, L. Allegra, Ed., Proceedings of a Conference in conjunction with the 1988 World Materials Congress (Chicago), Sept 1988, ASM INTERNATIONAL, 1988 9. Y. Shindou et al., Properties of Organic Composite-Coated Steel Sheet for Automobile Body Panels (SAE Paper 862016), in Proceedings of the Automotive Corrosion and Prevention Conference, P-188, Society of Automotive Engineers, 1986, p 81−90 10. M. Yamashita, T. Kubota, and T. Adaniya, Organic-Silicate Composite Coated Steel Sheet for Automobile Body Panel (SAE Paper 862017), in Proceedings of the Automotive Corrosion and Prevention Conference, P-188, Society of Automotive Engineers, 1986, p 91-97 11. T. Mohri et al., Newly Developed Organic Composite-Coated Steel Sheet With Bake Hardenability (SAE Paper 862030), in Proceedings of the Automotive Corrosion and Prevention Conference, P-188, Society of Automotive Engineers, 1986, p 199-208 12. T.E. Dorsett, Development of a Composite Coating for Pre-Coated Automotive Sheet Metal (SAE Paper 862027), in Proceedings of the Automotive Corrosion and Prevention Conference, P-188, Society of Automotive Engineers, 1986, p 163-173 13. T. Watanabe, T. Kubota, M. Yamashita, T. Urakawa, and M. Sagiyama, Corrosion-Resistant Precoated Steel Sheets for Automotive Body Panels, in Corrosion-Resistant Automotive Sheet Steels, L. Allegra, Ed., Proceedings of a Conference in conjunction with the 1988 World Materials Congress (Chicago), Sept 1988, ASM INTERNATIONAL, 1988


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Carbon and Low-Alloy Steel Plate

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Carbon and Low-Alloy Steels Carbon and Low-Alloy Steel Plate Revised by F.B. Fletcher, Lukens Steel Company STEEL PLATE is any flat-rolled steel product more than 200 mm (8 in.) wide and more than 6.0 mm (0.230 in.) thick or more than 1220 mm (48 in.) wide and 4.6 mm (0.180 in.) thick. The majority of mills for rolling steel plate have a working-roll width between 2030 and 5600 mm (80 and 220 in.). Therefore, the width of product normally available ranges from 1520 to 5080 mm (60 to 200 in.). Most steel plate consumed in North America ranges in width from 2030 to 3050 mm (80 to 120 in.) and ranges in thickness from 5 to 200 mm (3=16to 8 in.). Some plate mills, however, have the capability to roll steel more than 640 mm (25 in.) thick. Steel plate is usually used in the hot-finished condition, but the final rolling temperature can be controlled to improve both strength and toughness. Heat treatment is also used to improve the mechanical properties of some plate. Steel plate is mainly used in the construction of buildings, bridges, ships, railroad cars, storage tanks, pressure vessels, pipe, large machines, and other heavy structures, for which good formability, weldability, and machinability are required. The impairment of these desirable characteristics with increasing carbon content usually limits the steel to the low-carbon and medium-carbon constructional grades, with the low-carbon grades predominating. Many alloy steels are also produced as plate. In the final structure, however, alloy steel plate is sometimes heat treated to achieve mechanical properties superior to those typical of the hot-finished product.

Steelmaking Practices Steel plate is produced from continuously cast slabs or individually cast ingots or slabs. Preparing these steel slabs or ingots for subsequent forming into plates may involve requirements regarding deoxidation practices, austenite grain size, and/or secondary melting practices. Deoxidation Practices. During the steelmaking process, segregation of carbon can occur when carbon reacts with the dissolved oxygen in the molten steel (a reaction that is favored thermodynamically at lower temperatures). Therefore, the practice of controlling dissolved oxygen in the molten metal before and during casting is an important factor in improving the internal soundness and chemical homogeneity of cast steel. Deoxidation is also important in lowering the impact transition temperatures. Deoxidation can be achieved by vacuum processing or by adding deoxidizing elements such as aluminum or silicon. Steels are classified by their level of deoxidation: killed steel, semikilled steel, capped steel, and rimmed steel. The steel used for plates is usually either killed or semikilled. Semikilled steel is commonly used for casting ingots because it is more economical than killed steel. Continuously cast steels are normally fully killed to assure internal soundness. Killed steel is fully deoxidized, and from the viewpoint of minimum chemical segregation and uniform mechanical properties, killed steel represents the best quality available. Therefore, killed steel is generally specified when homogeneous structure and internal soundness of the plate are required or when improved low-temperature impact properties are desired. Killed steel can be produced either fine or coarse grained without adversely affecting soundness, surface, or cleanliness. Generally, heavy-gage plate (thicker than 38 mm, or 11=2in.) is produced from killed steel to provide improved internal homogeneity. Semikilled steel is deoxidized to a lesser extent that killed steel and therefore does not have the same degree of chemical uniformity or freedom from surface imperfections as killed steel. This type of steel is used primarily on lighter-gage plate, for which high reductions from ingot to plate thicknesses minimize the structural and chemical variations found in the as-cast ingot. Austenitic Grain Size. Steel plate specifications for structural and pressure vessel applications may require a steelmaking process that produces a fine austenitic grain size. When a fine austenitic grain size is specified, grain-refining elements are added during steelmaking. Aluminum is effective in retarding austenitic grain growth, resulting in improved toughness for heat-treated (normalized or quenched and tempered) steels. Steels used in high-temperature service normally contain only very small quantities of aluminum because aluminum may affect strain-aging characteristics and graphitization. However, the addition of aluminum may be necessary for some high-temperature steels (as well as most low-temperature steels) requiring good toughness. Other grain-refining elements, such as niobium, vanadium, and titanium, are used in high-strength low-alloy (HSLA) steels for grain refinement during rolling (see the article "High-Strength Structural and High-Strength Low-Alloy Steels" in this Volume). Melting Practices. The steel for plate products can be produced by the following primary steelmaking processes: open hearth, basic oxygen, or electric furnace. In addition, the steel can be further refined by secondary processes such as vacuum degassing or various ladle treatments for deoxidation or desulfurization.


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Vacuum degassing is used to remove dissolved oxygen and hydrogen from steel, thus reducing the number and size of indigenous nonmetallic inclusions. It also reduces the likelihood of internal fissures or flakes caused when hydrogen content is higher than desired. A small cost premium is associated with the specification of vacuum degassing. Desulfurization. By combining steel refining with the addition of ladle desulfurizing agents (for example, calcium or rare earth additions) immediately before casting or teeming, final plate steel sulfur content can be reduced to less than 0.005%. Lower sulfur content improves plate through-thickness properties and impact properties, but adds to the cost of the steel.

Platemaking Practices As noted earlier, steel plates are produced from either continuous-cast slabs, pressure-cast slabs, or ingots. Steel ingots are typically between 380 and 1140 mm (15 and 45 in.) thick. There ingots first pass through a slabbing mill where they are reduced in thickness to make a slab. The slab is then inspected, and the surface is conditioned by grinding or scarfing to remove surface imperfections, and then reheated in furnaces prior to rolling to final plate thickness. Continuous-cast slabs and pressure-cast slabs are normally heated and rolled to final plate thickness in a single operation. The plate can then be roller leveled and cooled. Microalloyed HSLA steels can be controlled rolled for grain refinement (see the article "High-Strength Structural and High-Strength Low-Alloy Steels" in thisVolume). In this case, the reheating temperature is lower than usual, and the rolling practices are designed to impart heavy reductions at relatively low temperatures. This form of thermomechanically controlled processing (TMCP) is used for grain refinement, which results in plates with improved toughness and strength compared to conventional plate rolling. In some plate mills, controlled rolling is followed by accelerated cooling or direct quenching instead of air cooling. Attractive combinations of strength and toughness can be achieved by TMCP. After cooling, plates are cut to size by shearing or thermal cutting. Following this operation, testing to confirm mechanical properties is customarily performed, and then the material is shipped to the fabricator. Certain plate products, however, require further processing such as heat treatment.

Plate Imperfections Certain characteristic surface imperfections that can weaken the plate may appear on hot-finished steel; chemical segregation that can alter properties across the section may also be present. Some of these imperfections are discussed below. Seams are the most common imperfections found in hot-finished steel. These longitudinal cracks on the surface are caused by blowholes and cracks in the original ingot that have been rolled closed, but not welded. For many plate applications, seams are of minor consequence. However, seams are harmful for applications involving heat treating or upsetting or in certain parts subjected to fatigue loading. Decarburization, a surface condition common to all hot-finished steel, is produced during the heating and rolling operations when atmospheric oxygen reacts with the heated surface, removing carbon. This produces a soft, low-strength surface, which is often unsatisfactory for applications involving wear or fatigue. For this reason, critical parts or at least critical areas of parts are usually machined to remove this weakened surface. Segregation. Alloying elements always segregation during the solidification of steel. Elements that are especially prone to segregation are carbon, phosphorus, sulfur, silicon, and manganese. The effect of segregation on mechanical properties and fabricability is insignificant for most plate steel applications. However, segregation may produce difficulties in subsequent operations such as forming, welding, punching, and machining.

Heat Treatment Although most steel plate is used in the hot-finished condition, the following heat treatments are applied to plate that must meet special requirements. Normalizing consists of heating the steel above its critical temperature and cooling in air. This refines the grain size and provides improved uniformity of structure and properties of the hot-finished plate. When toughness requirements are specified for certain thicknesses in some grades of normalized plate, accelerated cooling must be used in lieu of cooling in still air from the normalizing temperature. Such cooling is accomplished by fans to provide air circulation during cooling or by a water spray or dip. Accelerated cooling is used most often in plates with heavy thicknesses to obtain properties comparable to those developed by normalizing material in the lighter thicknesses. Quenching consists of heating the steel to a suitable austenitizing temperature, holding at that temperature, and quenching in a suitable medium that depends on chemical composition and cross-sectional dimensions. As-quenched steels are hard, high in strength, and brittle. They are almost always tempered before being placed in service. Tempering consists of reheating the steel to a predetermined temperature below the critical range, then cooling under suitable conditions. This treatment is usually carried out after normalizing or quenching to obtain desired mechanical properties. Those include a balance of strength and toughness to meet the designer's requirements. Stress relieving consists of heating the steel to a subcritical temperature to release stresses induced by such operations as flattening or other cold working, shearing, or gas cutting. Stress relieving is not intended to significantly modify the microstructure or to obtain desired mechanical properties.

Types of Steel Plate


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Steel plate is classified according to composition, mechanical properties, and steel quality. The three general categories of steel plate considered in this article are carbon steel plate, low-alloy plate, and high-strength low-alloy (HSLA) steel plate. These three categories of steel plate are available in the steel plate quality levels given in Table 1 . Further discussion on these various quality levels is provided in the section "Steel Plate Quality" in this article. Table 1 Available quality levels for carbon, HSLA, and low-alloy steel plate Carbon steel plates

HSLA steel plates

Low-alloy steel plates

...

...

Regular quality Structural quality

Structural quality

Structural quality

Drawing quality

Drawing quality

Drawing quality

Cold-drawing quality



Cold-pressing quality

...

Cold-flanging quality

Cold-flanging quality

Forging quality

Forging quality

Pressure vessel quality

Pressure vessel quality ...

... Cold-flanging quality ... Pressure vessel quality

...

Aircraft quality

Source: Ref 1

General Categories Carbon steel plate is available in all quality levels except aircraft quality (Table 1 ) and is available in many grades. Generally, carbon steel contains carbon up to about 2% and only residual quantities of other elements except those added for deoxidation, with silicon usually limited to 0.60% and manganese to about 1.65%. The chemical composition requirements of standard carbon steel plate are listed in Table 2 . These steels may be suitable for some structural applications when furnished according to ASTM A 830 and A 6. In addition to the carbon steels listed in Table 2 , other carbon steel plates are also classified according to more specific requirements in various ASTM specifications (see the section "Steel Plate Quality" in this article). Table 2 Standard carbon steel plate compositions applicable for structural applications When silicon is required, the following ranges and limits are commonly used for nonresulfurized carbon steel: 0.10% max, 0.07−0.15%, 0.10−0.20%, 0.15−0.30%, 0.35% max, 0.20−0.40, or 0.30−0.60%. Steel designation Chemical composition limits, % SAE or AISI No.

C

Mn

P(a)

S(a)

G10060

1006

0.08(a)

0.45(a)

0.040

0.050

G10080

1008

0.10(a)

0.50(a)

0.040

0.050

G10090

1009

0.15(a)

0.60(a)

0.040

0.050

G10100

1010

0.08−0.13

0.30−0.60

0.040

0.050

G10120

1012

0.10−0.15

0.30−0.60

0.040

0.050

G10150

1015

0.12−0.18

0.30−0.60

0.040

0.050

G10160

1016

0.12−0.18

0.60−0.90

0.040

0.050

G10170

1017

0.14−0.20

0.30−0.60

0.040

0.050

G10180

1018

0.14−0.20

0.60−0.90

0.040

0.050

G10190

1019

0.14−0.20

0.70−1.00

0.040

0.050

G10200

1020

0.17−0.23

0.30−0.60

0.040

0.050

G10210

1021

0.17−0.23

0.60−0.90

0.040

0.050

G10220

1022

0.17−0.23

0.70−1.00

0.040

0.050

G10230

1023

0.19−0.25

0.30−0.60

0.040

0.050

G10250

1025

0.22−0.28

0.30−0.60

0.040

0.050

G10260

1026

0.22−0.28

0.60−0.90

0.040

0.050

G10300

1030

0.27−0.34

0.60−0.90

0.040

0.050

G10330

1033

0.29−0.36

0.70−1.00

0.040

0.050

UNS


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01 Sep 2005

G10350

1035

0.31−0.38

0.60−0.90

0.040

0.050

G10370

1037

0.31−0.38

0.70−1.00

0.040

0.050

G10380

1038

0.34−0.42

0.60−0.90

0.040

0.050

G10390

1039

0.36−0.44

0.70−1.00

0.040

0.050

G10400

1040

0.36−0.44

0.60−0.90

0.040

0.050

G10420

1043

0.39−0.47

0.60−0.90

0.040

0.050

G10430

1043

0.39−0.47

0.70−1.00

0.040

0.050

G10450

1045

0.42−0.50

0.60−0.90

0.040

0.050

G10460

1046

0.42−0.50

0.70−1.00

0.040

0.050

G10490

1049

0.45−0.53

0.60−0.90

0.040

0.050

G10500

1050

0.47−0.55

0.60−0.90

0.040

0.050

G10550

1055

0.52−0.60

0.60−0.90

0.040

0.050

G10600

1060

0.55−0.66

0.60−0.90

0.040

0.050

G10640

1064

0.59−0.70

0.50−0.80

0.040

0.050

G10650

1065

0.59−0.70

0.60−0.90

0.040

0.050

G10700

1070

0.65−0.76

0.60−0.90

0.040

0.050

G10740

1074

0.69−0.80

0.50−0.80

0.040

0.050

G10750

1075

0.69−0.80

0.40−0.70

0.040

0.050

G10780

1078

0.72−0.86

0.30−0.60

0.040

0.050

G10800

1080

0.74−0.88

0.60−0.90

0.040

0.050

G10840

1084

0.80−0.94

0.60−0.90

0.040

0.050

G10850

1085

0.80−0.94

0.70−1.00

0.040

0.050

G10860

1086

0.80−0.94

0.30−0.50

0.040

0.050

G10900

1090

0.84−0.98

0.60−0.90

0.040

0.050

G10950

1095

0.90−1.04

0.30−0.50

0.040

0.050

G15240

1524

0.18−0.25

1.30−1.65

0.040

0.050

G15270

1527

0.22−0.29

1.20−1.55

0.040

0.050

G15360

1536

0.30−0.38

1.20−1.55

0.040

0.050

G15410

1541

0.36−0.45

1.30−1.65

0.040

0.050

G15480

1548

0.43−0.52

1.05−1.40

0.040

0.050

1552

0.46−0.55

1.20−1.55

0.040

0.050

G15520 (a) Maximum

Low-Alloy Steel Plate. Steel is considered to be low-alloy steel when either of the following conditions is met: • The maximum of the range given for the content of alloying elements exceeds one or more of the following limits: 1.65% Mn, 0.60% Si, and 0.60% Cu • Any definite range or definite minimum quantity of any of the following elements is specified or required within the limits of the recognized field of constructional alloy steels: aluminum, boron, chromium up to 3.99%, cobalt, niobium, molybdenum, nickel, titanium, tungsten, vanadium, zirconium, or any other alloying element added to obtain the desired alloying effect Alloying elements are added to hot-finished plates for various reasons, including improved corrosion resistance and/or improved mechanical properties at low or elevated temperatures. Alloying elements are also used to improve the hardenability of quenched and tempered plate. Low-alloy steels generally require additional care throughout their manufacture. They are more sensitive to thermal and mechanical operations, the control of which is complicated by the varying effects of different chemical compositions. To secure the most satisfactory results, consumers normally consult with steel producers regarding the working, machining, heat treating, or other operations to be employed in fabricating the steel; mechanical operations to be employed in fabricating the steel; mechanical properties to be obtained; and the conditions of service for which the finished articles are intended. The chemical composition requirements of standard low-alloy steel plate are listed in Table 3 . These low-alloy steels may be suitable for some structural applications when furnished according to ASTM A 6 and A 829. The effect of residual alloying elements on the mechanical properties of hot-finished steel plate is discussed in the section "Mechanical Properties" in this


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article. The effect of alloying elements on the hardenability and mechanical properties of quenched and tempered steels is discussed in the articles "Hardenable Carbon and Low-Alloy Steels" and "High-Strength Structural and High-Strength Low-Alloy Steels" in this Volume. Table 3 Composition ranges and limits for AISI-SAE standard low-alloy steel plate applicable for structural applications Boron or lead can be added to these compositions. Small quantities of certain elements not required may be found. These elements are to be considered incidental and are acceptable to the following maximum amounts: copper to 0.35%, nickel to 0.25%, chromium to 0.20%, and molybdenum to 0.06%. Heat composition ranges and limits, %(a) AISI-SAE UNS designation

designation

C

Mn

Si(b)

Cr

Ni

Mo

1330

G13300

0.27−0.34

1.50−1.90

0.15−0.30

...

...

...

1335

G13350

0.32−0.39

1.50−1.90

0.15−0.30

...

...

...

1340

G13400

0.36−0.44

1.50−1.90

0.15−0.30

...

...

...

1345

G13450

0.41−0.49

1.50−1.90

0.15−0.30

...

...

...

4118

G41180

0.17−0.23

0.60−0.90

0.15−0.30

0.40−0.65

...

0.08−0.15

4130

G41300

0.27−0.34

0.35−0.60

0.15−0.30

0.80−1.15

...

0.15−0.25

4135

G41350

0.32−0.39

0.65−0.95

0.15−0.30

0.80−1.15

...

0.15−0.25

4137

G41370

0.33−0.40

0.65−0.95

0.15−0.30

0.80−1.15

...

0.15−0.25

4140

G41400

0.36−0.44

0.70−1.00

0.15−0.30

0.80−1.15

...

0.15−0.25

4142

G41420

0.38−0.46

0.70−1.00

0.15−0.30

0.80−1.15

...

0.15−0.25

4145

G41450

0.41−0.49

0.70−1.00

0.15−0.30

0.80−1.15

...

0.15−0.25

4340

G43400

0.36−0.44

0.55−0.80

0.15−0.30

0.60−0.90

1.65−2.00

0.20−0.30

E4340(c)

G43406

0.37−0.44

0.60−0.85

0.15−0.30

0.65−0.90

1.65−2.00

0.20−0.30

4615

G46150

0.12−0.18

0.40−0.65

0.15−0.30

...

1.65−2.00

0.20−0.30

4617

G46170

0.15−0.21

0.40−0.65

0.15−0.30

...

1.65−2.00

0.20−0.30

4620

G46200

0.15−0.30

...

1.65−2.00

0.20−0.30

5160

G51600

0.54−0.65

0.70−1.00

0.15−0.30

0.60−0.90

...

...

6150(d)

G61500

0.46−0.54

0.60−0.90

0.15−0.30

0.80−1.15

...

...

8615

G86150

0.12−0.18

0.60−0.90

0.15−0.30

0.35−0.60

0.40−0.70

0.15−0.25

8617

G86170

0.15−0.21

0.60−0.90

0.15−0.30

0.35−0.60

0.40−0.70

0.15−0.25

8620

G86200

0.17−0.23

0.60−0.90

0.15−0.30

0.35−0.60

0.40−0.70

0.15−0.25

8622

G86220

0.19−0.25

0.60−0.90

0.15−0.30

0.35−0.60

0.40−0.70

0.15−0.25

8625

G86250

0.22−0.29

0.60−0.90

0.15−0.30

0.35−0.60

0.40−0.70

0.15−0.25

8627

G86270

0.24−0.31

0.60−0.90

0.15−0.30

0.35−0.60

0.40−0.70

0.15−0.25

8630

G86300

0.27−0.34

0.60−0.90

0.15−0.30

0.35−0.60

0.40−0.70

0.15−0.25

8637

G86370

0.33−0.40

0.70−1.00

0.15−0.30

0.35−0.60

0.40−0.70

0.15−0.25

8640

G86400

0.36−0.44

0.70−1.00

0.15−0.30

0.35−0.60

0.40−0.70

0.15−0.25

8655

G86550

0.49−0.60

0.70−1.00

0.15−0.30

0.35−0.60

0.40−0.70

0.15−0.25

0.16−0.22

0.40−0.65

8742 G87420 0.38−0.46 0.70−1.00 0.15−0.30 0.35−0.60 0.40−0.70 0.20−0.30 (a) Indicated ranges and limits apply to steels made by the open hearth or basic oxygen processes; maximum content for phosphorus is 0.035% and for sulfur 0.040%. For steels made by the electric furnace process, the ranges and limits are reduced as follows: C⎯0.01%; Mn⎯0.05%; Cr⎯0.05% (1.25%); maximum content for either phosphorus or sulfur is 0.025%. (b) Other silicon ranges may be negotiated. Silicon is available in ranges of 0.10−0.20%, 0.20−0.30%, and 0.35% maximum (when carbon deoxidized) when so specified by the purchaser. (c) Prefix "E" indicates that the steel is made by the electric furnace process. (d) Contains 0.15% V minimum

In addition to the low-alloy steels listed in Table 3 , other low-alloy steel plates are also classified according to more specific requirements in various ASTM specifications. The chemical composition requirements and mechanical properties of low-alloy steel plate in ASTM standards are discussed in the section "Steel Plate Quality" in this article. High-strength low-alloy steels offer higher mechanical properties and, in certain of these steels, greater resistance to atmospheric corrosion than conventional carbon structural steels. The HSLA steels are generally produced with emphasis on mechanical property requirements rather than the chemical composition limits. They are not considered alloy steels as described in the American Iron and Steel Institute (AISI) steel products manuals, even though utilization of any intentionally added alloy


Page 379



01 Sep 2005

content would technically qualify as such. There are two groups of compositions in this category: • Vanadium and/or niobium steels, with a manganese content generally not exceeding 1.35% maximum and with the addition of 0.2% minimum copper when specified • High-strength intermediate-manganese steels, with a manganese content in the range of 1.10 to 1.65% and with the addition of 0.2% minimum copper when specified Other elements commonly added to HSLA steels to yield the desired properties include silicon, chromium, nickel, molybdenum, titanium, zirconium, boron, aluminum, and nitrogen. The chemical compositions of ASTM structural quality and pressure vessel quality plates made of HSLA steel are listed in Table 4 . More information on HSLA steels is provided in the article "High-Strength Structural and High-Strength Low-Alloy Steels" in this Volume. Table 4 Composition of high-strength low-alloy steel plate ASTM Material specificati grade on or type

Composition, %(a) C

Mn

P

S

Si

Cr

Ni

Mo

Cu

V

Nb

Others

0.25

0.40

0.08

0.35

0.10

0.05

...

Structural quality A 131

AH32, DH32, EH32, AH36, DH36, EH36

0.18

A 242

1

0.15

1.00

0.15

0.05

...

...

...

...

0.20 min

...

...

(b)(c)

A 572

42

0.21

1.35

0.04

0.05

0.40(d)

...

...

...

...

(e)

(e)

(e)

45

0.22

1.35

0.04

0.05

0.40(d)

...

...

...

...

(e)

(e)

(e)

50

0.23

1.35

0.04

0.05

0.40(d)

...

...

...

...

(e)

(e)

(e)

60

0.26

1.35

0.04

0.05

0.40(d)

...

...

...

...

(e)

(e)

(e)

0.04

0.05

0.40

...

...

(e)

65 A 588

A 656

A 678

0.26(d) 1.65(d)

0.04 0.10−0.50

...

...

(e)

(e)

A

0.19

0.80−1.2 0.04 5

0.05 0.30−0.65 0.40−0 .65

0.40

...

0.25−0.4 0.02−0.10 0

...

...

B

0.20

0.75−1.3 0.04 5

0.05 0.15−0.50 0.50−0 .70

0.50

...

0.20−0.4 0.01−0.10 0

...

...

C

0.15

0.80−1.3 0.04 5

0.05 0.15−0.40 0.30−0 0.25−0 .50 .50

...

0.20−0.5 0.01−0.10 0

...

...

0.05 0.50−0.90 0.50−0 .90

...

...

...

0.04

Zr, 0.05−0.15

0.40

...

0.005−0.0 5(f)

...

D

A 633

0.90−1.6 0.04 0

0.10−0 0.75−1.2 0.04 .20 5

0.30

K

0.17

0.50−1.2 0.04 0

0.05 0.25−0.50 0.40−0 .70

A

0.18

1.00−1.3 0.04 5

0.05 0.15−0.50

...

...

...

...

...

0.05

...

B

0.18

1.00−1.3 0.04 5

0.05 0.15−0.50

...

...

...

...

0.10

...

...

C

0.20

1.15−1.5 0.04 0

0.05 0.15−0.50

...

...

...

...

...

0.01−0.05

...

D

0.20

0.70−1.6 0.04 0(d)

0.05 0.15−0.50

0.25

0.25

0.08

0.35

...

...

...

E

0.22

1.15−1.5 0.04 0

0.05 0.15−0.50

...

...

...

...

0.04−0.11

(g)

N, 0.01−0.03(h )

3

0.18

1.65

0.025 0.035

0.60

...

...

...

...

0.08

0.005−0.1 5

N, 0.020

7

0.18

1.65

0.025 0.035

0.60

...

...

...

...

0.005−0.0 0.005−0.0 15(i) 15(i)

N,0.020

D

0.22

...

...

...

0.2 min(j)

1.15−1.5 0.04 0

0.05 0.15−0.50


0.10 0.30−0.5 0

0.04−0.11

(g)

N, 0.001−0.03

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A 709

50

0.23


1.359(d) 0.04

0.05 0.15−0.40 (d)

50W

...

...

...

01 Sep 2005

...

(e)

(e)

(e)

0.10

0.02−0.10

(Nb + V), 0.15

...

...

Identical to A 588 type A, B, or C (as specified)

A 808

...

0.12

1.65

0.04

A 852

...

0.19

...

...

0.05 0.20−0.65 0.40−0 .70

0.50

...

A 871

...

1.20

1.50

0.04

0.05

0.90

0.90

1.25

0.25

1.00

0.10

0.05

Zr, 0.15; Ti, 0.05

1.60

0.035 0.015

0.40

...

0.35

0.25

0.25(j)

0.11

(k)

Al, 0.06; N, 0.030

0.80−1.3 0.04 5

0.05 0.15−0.50

...

...

0.20−0.4 0.02−0.10 0

Pressure vessel quality A 734

B

0.17

A 737

B

0.20

1.15−1.5 0.035 0.030 0.15−0.50 0

...

...

...

...

...

0.05

...

C

0.22

1.15−1.5 0.035 0.030 0.15−0.50 0

...

...

...

...

0.04−0.11

(k)

N, 0.03

A 841

...

0.20

0.25 0.08 0.35 0.06 0.03 Al, 0.020 0.70−1.6 0.030 0.030 0.15−0.50 0.25 min 0(d) (a) Except as noted, when a single value is shown, it is a maximum limit. (b) Choice and amount of other alloying elements added to give the required mechanical properties and atmospheric corrosion resistance are made by the producer and reported in the heat analysis. (c) Elements commonly added include silicon, chromium, nickel, vanadium, titanium, and zirconium. (d) Limiting values vary with plate thickness. (e) For type 1, 0.005−0.05% Nb; for type 2, 0.01−0.15% V; for type 3, 0.05% Nb max + V = (0.02−0.15%); for type 4, N (with V) 0.015% max. (f) For plates under 13 mm (1=2in.) thickness, the minimum niobium limit is waived. (g) Niobium may be present in the amount of 0.01−0.05%. (h) The minimum total aluminum content shall be 0.018% or the vanadium:nitrogen ratio shall be 4:1 minimum. (i) Niobium, or vanadium, or both, 0.005% min. When both are added, the total shall be 0.20% max. (j) Applicable only when specified. (k) 0.05% max Nb may be present.

Steel Plate Quality Steel quality, as the term applies to steel plate, is indicative of many conditions, such as the degree of internal soundness, relative uniformity of mechanical properties and chemical composition, and relative freedom from injurious surface imperfections. The various types of steel plate quality are indicated in Table 1 . The three main quality descriptors used to describe steel plate are regular quality, structural quality, and pressure vessel quality. Special qualities include cold-drawing quality, cold-pressing quality, cold-flanging quality, and forging quality carbon steel plate, along with drawing quality and aircraft quality alloy steel plate. Quality descriptors that have been used in the past include flange quality and firebox quality carbon and alloy steel plate and marine quality carbon steel plate. However, use of these descriptors has been discontinued in favor of pressure vessel quality. Regular quality is the most common quality of carbon steel, which is applicable to plates with a maximum carbon content of 0.33%. Plates of this quality are not expected to have the same degree of chemical uniformity, internal soundness, or freedom from surface imperfections that is associated with structural quality or pressure vessel quality plate. Regular quality is usually ordered to standard composition ranges and is not customarily produced to mechanical property requirements. Regular quality is analogous to merchant quality for bars because there are normally no restrictions on deoxidation, grain size, check analysis, or other metallurgical factors. Also, this quality plate can be satisfactorily used for applications similar to those of merchant quality bars, such as those involving mild cold bending, mild hot forming, punching, and welding for noncritical parts of machinery. Structural quality steel plate is intended for general structural applications such as bridges, buildings, transportation equipment, and machined parts. The various ASTM specifications for structural quality steel plate are given in Table 5 . Most of the structural steel plate listed in Table 5 is furnished to both chemical composition limits (Table 6 ) and mechanical properties (Table 7 ). However, some structural steel plate (ASTM A 829 and A 830 in Table 5 ) is produced from the standard steels listed in Tables 2 and 3 . These steels can be furnished only according to the chemical compositions specified by SAE/AISI steel designations. Factors affecting the mechanical properties of hot-finished carbon steel are discussed in the section "Mechanical Properties" in this article. Table 5 ASTM specifications for structural quality steel plate General requirements for structural plate are covered in ASTM A 6. ASTM specification(a)

Steel type and condition

Carbon steel A 36(b)

Carbon steel shapes, plates, and bars of structural quality

A 131(c)

Structural steel shapes, plates, bars, and rivets for use in ship construction (ordinary strength)

A 283(b)

Low and intermediate tensile strength carbon steel plates


Page 381



A 284

01 Sep 2005

Low and intermediate tensile strength carbon-silicon steel plates for machine parts and general construction

A 529(d)

Structural steel with 290 MPa (42 ksi) minimum yield point

A 573

Structural quality carbon-manganese-silicon steel plates with improved toughness

A 678

Quenched and tempered carbon and HSLA plates for structural applications

A 709

Carbon and HSLA steel structural shapes, plates, and bars, and quenched and tempered alloy steel for use in bridges

A 827(e)

Carbon steel plates for forging applications

A 830(e)

Structural quality carbon steel plates furnished to chemical requirements

Low-alloy steel A 514

Structural quality quenched and tempered alloy steel plates for use in welded bridges and other structures

A 709

See above under "Carbon steel"

A 710

Low-carbon age-hardening Ni-Cu-Cr-Mo-Nb, Ni-Cu-Nb, and Ni-Cu-Mn-No-Nb alloy steel plates, shapes, and bars for general applications.

A 829(e)(f)

Structural quality alloy plates specified to chemical composition requirements

HSLA steel A 13(c)

Structural steel shapes, plates, bars, and rivets for use in ship construction (higher strength)

A 242

HSLA structural steel shapes, plates, and bars for welded, riveted, or bolted construction

A 441(g)

Mn-V HSLA steel plates, bars, and shapes

A 572

HSLA structural Nb-V steel shapes, plates, sheet piling, and bars for riveted, bolted, or welded construction of bridges, buildings, and other structures

A 588(h)

HSLA structural steel shapes, plates, and bars for welded, riveted, or bolted construction for use in bridges and buildings with atmospheric corrosion resistance approximately two times that of carbon steel with copper

A 633

Normalized HSLA structural steel for welded, riveted, or bolted construction suited for service at low ambient temperatures of −45 °C (−50 °F) or higher

A 656

Hot-rolled HSLA structural steel with improved formability for use in truck frames, brackets, crane booms, rail cars, and similar applications

A 678


A 709


A 808

Hot-rolled HSLA Mn-V-Nb structural steel plate with improved notch toughness

A 852

Quenched and tempered HSLA structural steel plate for welded, riveted, or bolted construction for use in bridges and buildings with atmospheric corrosion resistance approximately two times that of carbon steel with copper

A 871

HSLA structural steel plate in the as-rolled, normalized, or quenched and tempered condition with atmospheric corrosion resistance approximately two times that of carbon steel with copper (a) Also designated with the suffix "M" when the specification covers metric equivalents. (b) This specification is also published by the American Society of Mechanical Engineers, which uses the prefix "S" (for example, SA36). (c) See also Section 43 of the American Bureau of Shipping specifications and MIL-S-22698 (SH). (d) 13 mm (1=2in.) maximum thickness. (e) See also Ref 1. (f) Tensile properties may also be specified when compatible. (g) Discontinued in 1989 and replaced by A 572. (h) Minimum yield point 345 MPa (50 ksi) to 100 mm (4 in.). Lower minimum yield points for thicker sections.

Table 6 ASTM specifications of chemical composition for structural plate made of low-alloy steel or carbon steel Mate rial grade ASTM or specificati type on

Composition, %(a)

C

Mn

P

S

Si

Cr

Ni

Mo

Cu

0.15−0.21

0.80−1.10

0.035

0.04

0.40−0.80

0.50−0.80

...

0.18−0.28

...

Others


A


Zr, 0.05−0.15; B, 0.0025

Page 382



01 Sep 2005

B

0.12−0.21

0.70−1.00

0.035

0.04

0.20−0.35

0.40−0.65

...

0.15−0.25

...

V, 0.03−0.08; Ti, 0.01−0.03; B, 0.0005−0.005

C

0.10−0.20

1.10−1.50

0.035

0.04

0.15−0.30

...

...

0.15−0.30

...

B, 0.001−0.005

E

0.12−0.20

0.40−0.70

0.035

0.04

0.20−0.40

1.40−2.00

...

0.40−0.60

...

Ti, 0.01−0.10 (b), 0.001−0.005

F

0.10−0.20 0.0.60−1.0 0

0.035

0.04

0.15−0.35

0.40−0.65 0.70−1.0 0.40−0.60 0

H

0.12−0.21

0.95−1.30

0.035

0.04

0.20−0.35

0.40−0.65 0.30−0.7 0.20−0.30 0

J

0.12−0.21

0.45−0.70

0.035

0.04

0.20−0.35

...

M

0.12−0.21

0.45−0.70

0.035

0.04

0.20−0.35

...

P

0.12−0.21

0.45−0.70

0.035

0.04

Q

0.14−0.21

0.95−1.30

0.035

R

0.15−0.80

0.85−1.15

...

V, 0.03−0.08; B, 0.0005−0.005

0.50−0.65

...

B, 0.001−0.005

1.20−1.5 0.45−0.60 0

...

B, 0.001−0.005

0.20−0.35

0.85−1.20 1.20−1.5 0.45−0.60 0

...

B, 0.001−0.005

0.04

0.15−0.35

1.00−1.50 1.20−1.5 0

0.40−0.6

...

V, 0.03−0.08

0.035

0.04

0.20−0.35

0.35−0.65

90−1.10

0.15−0.25

...

V, 0.03−0.08

0.04

...

0.10−0.35

...

B, 0.001−0.005; Nb, 0.06 max(c)

...

0.45−0.60

...

V, 0.03−0.08; B, 0.001−0.005

S

0.10−0.20

1.10−1.50

0.035

0.15−0.35

...

T

0.08−0.14

1.20−1.50

0.035 0.010 0.40−0.60

...

...

A 709

100, 100W

A 710

A

0.07

0.40−0.70

0.025 0.025

B

0.06

0.40−0.65

0.025 0.025 0.15−0.40

...

1.20−1.5 0

C

0.07

1.30−1.65

0.025

...

0.70−1.0 0.15−0.25 0

A 829

0.15−0.50 V, 0.03−0.08; B, 0.0005−0.006

(equivalent to A 514-A, B, C, E, F, H, J, M, P, Q)

0.25

0.40

0.60−0.90 0.70−1.0 0.15−0.25 0

0.04

(d)

...

1.00−1.30 Nb, 0.02 min 1.00−1.30 Nb, 0.02 min 1.00−1.30 Nb, 0.02 min

(See Table 3 .)

Carbon steel A 36

...

0.29(e)

A 131

A

0.26(e)

(g)

0.05

0.05

...

...

...

...

...

...

B

0.21

0.80−1.10( h)

0.04

0.04

0.35

...

...

...

...

...

D

0.21

0.70−1.35( e)(h)

0.04

0.04

0.10−0.35

...

...

...

...

...

E

0.18

0.70−1.35( h)

0.04

0.04

0.10−0.35

...

...

...

...

...

CS, DS

0.16

1.00−1.35( h)

0.04

0.04

0.10−0.35

...

...

...

...

...

A

0.14

0.90

0.04

0.05

0.04(e)

...

...

...

0.20(f)

...

B

0.17

0.90

0.04

0.05

0.04(e)

...

...

...

0.20(f)

...

A 283

A 284 A 529

0.80−1.20( e)

0.04

0.05 0.15−0.40( e)

...

...

...

0.20(f)

...

C

0.24

0.90

0.04

0.05

0.04(e)

...

...

...

0.20(f)

...

D

0.27

0.90

0.04

0.05

0.04(e)

...

...

...

0.20(f)

...

C

0.36(e)

0.90

0.04

0.05

0.15−0.40

...

...

...

...

...

D

0.35(e)

0.90

0.04

0.05

0.15−0.40

...

...

...

...

...

...

0.27

1.20

0.04

0.05

...

...

...

...

0.20(f)

...


Page 383


A 573


01 Sep 2005

58

0.23

0.60−0.90( h)

0.04

0.05

0.10−0.35

...

...

...

...

...

65

0.26(e)

0.85−1.20

0.04

0.05

0.15−0.40

...

...

...

...

...

70

0.28(e)

0.85−1.20

0.04

0.05

0.15−0.40

...

...

...

...

...

A

0.16

0.90−1.50

0.04

0.05

0.15−0.50

0.25

0.25

0.08

0.20(f)−0.3 5

...

B

0.20

0.70−1.60( e)

0.04

0.05

0.15−0.50

0.25

0.25

0.08

0.20(f)−0.3 5

...

C

0.22

1.00−1.60

0.04

0.05

0.20−0.50

0.25

0.25

0.08

0.20(f)−0.3 5

...

A 709

36

0.27(e)

0.80−1.20( e)

0.04

0.05 0.15−0.40( e)

...

...

...

...

...

A 827

(d)

A 678

(See Table 11 .)

A 830 (d) (See Table 2 .) (a) Note: See Table 4 for the compositions of structural plate made of HSLA steel. (b) When a single value is shown, it is a maximum limit, except for copper, for which a single value denotes a minimum limit. (c) Vanadium can be substituted for part or all of the titanium on a one-for-one basis. (d) Titanium may be present in levels up to 0.06% to protect the boron additions. (e) Specification covers many AISI/SAE grades and chemistries. (f) Limiting values vary with plate thickness. (g) Minimum value applicable only if copper-bearing steel is specified. (h) Plates over 13 mm (1=2in.) in thickness shall have a minimum manganese content not less than 2.5 times carbon content. (i) The upper limit of manganese may be exceeded provided C + 1/6 Mn does not exceed 0.40% based on heat analysis.

Table 7 ASTM specifications of mechanical properties for structural plate made of carbon steel, low-alloy steel, and HSLA steel Tensile strength(a) ASTM specification

Yield strength(a)

Minimum elongation( Minimum b) elongation(b) in 200 mm in 50 mm (2 in.), % (8 in.), %

Material grade or type

MPa

ksi

MPa

ksi

A 36

...

400−500

58−80

220−250(b)

32−36(b)

20

23

A 131

A, B, D, E, CS, DS

400−490

58−71

220(b)

32(b)

21(b)

24

A 283

A

310−415

45−60

165

24

27

30

B

345−405

50−65

185

27

25

28

C

380−485

55−70

205

30

22

25

Carbon steel

D

415−515(b)

60−75(b)

230

33

20

23

A 284

C

415

60

205

30

21

25

D

415

60

230

33

21

24

A 529

...

415−585

60−85

290

42

19

...

A 573

58

400−490

58−71

220

32

21

...

65

450−530

65−77

240

35

20

...

70

485−620

70−90

290

42

18

...

A

485−620

70−90

345

50

...

22

60

...

22

A 678

B

550−690

80−100

415

C

585−793(b)

85−115(b)

450(b)

65(b)

...

19

A 709

36

400−550

58−80

250

36

20

23

A 827(c)

(d)

(See the section "Forging Quality Plates " in this article.)

A 830(c)

(d)

(See text.)


All

690−895(b)

100−130(b)

620(b)

90(b)

...

16

A 709

100, 100W

700−915

100−130

635(b)

90(b)

...

15(c)

A 710

A (class 1)

585(b)

85(b)

515(b)

75(b)

...

20

A (class 2)

485(b)

70(b)

415(b)

60(b)

...

20


Page 384


A 829(c)


01 Sep 2005

A (class 3)

485(b)

70(b)

415(b)

60(b)

...

20

B

605(b)

88(b)

515(b)

75(b)

...

18

C (class 1)

690

100

620

90

20

C (class 3)

620(b)

90(b)

550(b)

80(b)

20

(d)

(See text.)

HSLA steels A 131

AH32, DH32, EH32

470−585

65−85

315

46

19

22

AH36, DH36, EH36

490−620

71−90

...

51

19

22

A 242

...

435(b)

63(b)

290(b)

42(b)

18

21

A 572

42

415

60

290

42

20

24

50

450

65

345

50

18

21

60

520

75

415

60

16

18

A 588 A 633

A 656

65

550

80

450

65

15

17

All

435(b)

63(b)

290(b)

42(b)

18

21

A

430−570

63−83

290

42

18

23

C, D

450−590(b)

65-85(b)

315(b)

46(b)

18

23

E

515−655(b)

75−95(b)

380(b)

55(b)

18

23

50

415

60

345

50

20

...

60

485

70

415

60

17

...

70

550

80

485

70

14

...

80

620

90

550

80

12

...

A 678

D

620−760

90−110

515

75

...

18

A 709

50

450

65

345

50

18

21

50W

485

70

345

50

18

21

...

415(b)

60(b)

290(b)

42(b)

18

22

A 852

...

620−760

90−110

485

70

...

19

A 871

60

520

75

415

60

16

18

A 808

65 550 80 450 65 15 17 (a) Where a single value is shown, it is a minimum. (b) Minimum and/or maximum values depend on plate width and/or thickness. (c) Specification does not specify mechanical properties. (d) Includes several AISI/SAE grades

Pressure Vessel Plate. Steel plate intended for fabrication into pressure vessels must conform to specifications different from those of similar plate intended for structural applications. The major differences between the two groups of specifications are that pressure vessel plate must meet requirements for notch toughness and has more stringent limits for allowable surface and edge imperfections. Table 8 lists the various ASTM specifications for pressure vessel steel plate. All of these steel plate specifications are furnished according to both chemical composition limits and mechanical properties. Table 8 ASTM specifications for pressure vessel quality steel plate General requirements for pressure vessel plate are covered in ASTM A 20 Specification


Carbon steel A 285(a)

Carbon steel plates of low or intermediate tensile strength

A 299(a)

Carbon-manganese-silicon steel plates

A 442(b)

Carbon steel plates for applications requiring low transition temperature

A 455(a)

Carbon-manganese steel plates of high tensile strength

A 515(a)

Carbon-silicon steel plates for intermediate-and higher-temperature service

A 516(a)

Carbon steel plates for moderate and lower-temperature service

A 537(a)

Heat-treated carbon-manganese-silicon steel plates

A 562

Titanium-bearing carbon steel plates for glass or diffused metallic coatings

A 612(a)

Carbon steel plates of high tensile strength for moderate-and lower-temperature service

A 662(a)

Carbon-manganese steel plates for moderate-and lower-temperature service


Page 385



01 Sep 2005

A 724

Quenched and tempered carbon steel plates for layered pressure vessels not subject to postweld heat treatment

A 738(a)

Heat-treated carbon manganese-silicon steel plates for moderate-and lower-temperature service

Low-alloy steel A 202(a)

Cr-Mn-Si alloy steel plates

A 203(a)

Nickel alloy steel plates

A 204(a)

Molybdenum alloy steel plates

A 225(a)

Mn-V alloy steel plates

A 302(a)

Mn-Mo and Mn-Mo-Ni alloy steel plates

A 353(a)

Double normalized and tempered 9% Ni alloy steel plates for cryogenic service

A 387(a)

Cr-Mo alloy steel plates for elevated-temperature service

A 517(a)

Quenched and tempered alloy steel plates of high tensile strength

A 533(a)

Quenched and tempered Mn-Mo and Mn-Mo-Ni alloy steel plates

A 542(a)

Quenched and tempered Cr-Mo alloy steel plates

A 543(a)

Quenched and tempered Ni-Cr-Mo alloy steel plates

A 553(a)

Quenched and tempered 8% and 9% Ni alloy steel plates

A 645(a)

Specially heat treated 5% Ni alloy steel plates for low-or cryogenic-temperature service

A 734

Quenched and tempered alloy and HSLA steel plates for low-temperature service

A 735

Low-carbon Mn-Mo-Nb alloy steel plates for moderate-and lower-temperature service

A 736

Age-hardening low-carbon Ni-Cu-Cr-Mo-Nb alloy steel plates

A 782

Quenched and tempered Mn-Cr-Mo-Si-Zr alloy pressure vessel steel plates

A 832

Cr-Mo-V-Ti-B alloy pressure vessel steel plates

A 844

9% Ni alloy pressure vessel steel plates produced by the direct-quenching process

HSLA steel A 734

See under "Alloy steel"

A 737(a)

HSLA steel plates for applications requiring high strength and toughness

A 841 Steel pressure vessel plate produced by the thermomechanical control processes (a) This specification is also published by the American Society of Mechanical Engineers, which adds an "S" in front of the "A" (for example, SA285). (b) Discontinued in 1991 (c)

The chemical composition limits of pressure vessel steel plate include a maximum phosphorus content of 0.035% and a maximum sulfur content of 0.040% by product analysis. The chemical compositions of various types of pressure vessel steel plate are given in Table 9 . Table 9 ASTM specification of chemical composition for pressure vessel plate made of carbon and low-alloy steel See Table 4 for the compositions of pressure vessel plate made of HSLA steel. The maximum limits per ASTM A 20 on unspecified elements are 0.40% Cu, 0.40% Ni, 0.30% Cr, 0.12% Mo, 0.03% V, and 0.02% Nb. Composition, % (a) Materi al grade ASTM or specificatio type n C Mn P S Si Cr Ni Mo Cu Others Carbon steel A 285

A

0.17

0.90

0.035 0.04

...

...

...

...

...

...

B

0.22

0.90

0.035 0.04

...

...

...

...

...

...

C

0.28

0.90

0.035 0.04

...

...

...

...

...

...

A 299

...

0.30(b)

0.90−1.50(b) 0.035 0.04

0.15−0.40

...

...

...

...

...

A 442

55

0.24(b)

0.80−1.10(b) 0.035 0.04

0.15−0.40

...

...

...

...

...

0.80−1.10(b) 0.035 0.04

60

0.27(b)

A 455

...

0.33

0.15−0.40

...

...

...

...

...

0.85−1.20

0.035 0.04

0.10

...

...

...

...

...

A 515

55

0.28(b)

0.90

0.035 0.04

0.15−0.40

...

...

...

...

...

60

0.31(b)

0.90

0.035 0.04

0.15−0.40

...

...

...

...

...


Page 386



01 Sep 2005

65

0.33(b)

0.90

0.035 0.04

0.15−0.40

...

...

...

...

...

70

0.35(b)

1.20

0.035 0.04

0.15−0.40

...

...

...

...

...

55

0.26(b)

0.60−1.20(b) 0.035 0.04

0.15−0.40

...

...

...

...

...

60

0.27(b)

0.60−1.20(b) 0.035 0.04

0.15−0.40

...

...

...

...

...

65

0.29(b)

0.85−1.20

0.035 0.04

0.15−0.40

...

...

...

...

...

70

0.31(b)

0.85−1.20

0.035 0.04

0.15−0.40

...

...

...

...

...

A 537

Class 1, 2

0.24

0.70−1.60(b) 0.035 0.04

0.15−0.50

0.25

0.25

0.08

0.35

...

A 562

...

0.12

0.15−0.40

...

...

...

A 612

...

0.29(b)

1.00−1.50(b) 0.035 0.04 0.15−0.50(b )

0.25

0.25

0.08

0.35

A 662

A

0.14

0.90−1.35

0.035 0.04

0.15−0.40

...

...

...

...

...

B

0.19

0.85−1.50

0.035 0.04

0.15−0.40

...

...

...

...

...

C

0.20

1.100−1.60

0.035 0.04

0.15−0.50

...

...

...

...

...

A

0.18

1.00−1.60

0.035 0.04

0.55

0.25

0.25

0.08

0.35

V, 0.08

B

0.20

1.00−1.60

0.035 0.04

0.50

0.25

0.25

0.08

0.35

V, 0.08

C

0.22

1.10−1.60

0.035 0.04

0.20−0.60

0.25

0.25

0.08

0.35

B, 0.005; V, 0.008

A

0.24

1.60(b)

0.035 0.04

0.15−0.50

0.25

0.50

0.08

0.35

B

0.20

0.90−1.50

0.030 0.025

0.15−0.55

0.25

0.25

0.08

0.35

V, 0.08

C

0.20

1.60(b)

0.030 0.025

0.15−0.55

0.25

0.25

0.08

0.35

V, 0.08

A

0.17

1.05−1.40

0.035 0.040

0.60−0.90

0.35−0.60

...

...

...

...

B

0.25

1.05−1.40

0.035 0.040

0.60−0.90

0.35−0.60

...

...

...

...

A

0.23(b)

0.80(b)

0.035 0.040

0.15−0.40

...

2.10−2.50

...

...

...

A 516

A 724

A 738

1.20

0.035 0.04

0.15 min Ti min, 4 × C V, 0.08

...

Low-alloy steel A 202 A 203

B

0.25(b)

0.80(b)

0.035 0.040

0.15−0.40

...

2.10−2.50

...

...

...

D

0.20(b)

0.80(b)

0.035 0.040

0.15−0.40

...

3.25−3.75

...

...

...

E,F

0.23(b)

0.80(b)

0.035 0.040

0.15−0.40

...

3.25−3.75

...

...

...

A

0.25(b)

0.90

0.035 0.040

0.15−0.40

...

...

0.45−0.60

...

...

B

0.27(b)

0.90

0.035 0.040

0.15−0.40

...

...

0.45−0.60

...

...

C

0.28(b)

0.90

0.035 0.040

0.15−0.40

...

...

0.45−0.60

...

...

C

0.25

1.60

0.035 0.040

0.15−0.40

...

0.40−0.70

...

...

V, 0.13−0.18

D

0.20

1.70

0.035 0.040

0.10−0.50

...

0.40−0.70

...

...

V, 0.10−0.18

A

0.25(b)

0.95−1.30

0.035 0.040

0.15−0.40

...

...

0.45−0.60

...

...

B

0.25(b)

1.15−1.50

0.035 0.040

0.15−0.40

...

...

0.45−0.60

...

...

C

0.25(b)

1.15−1.50

0.035 0.040

0.15−0.40

...

0.40−0.70 0.45−0.60

...

...

D

0.25(b)

1.15−1.50

0.035 0.040

0.15−0.40

...

0.70−1.00 0.45−0.60

...

...

A 353

...

0.13

0.90

0.035 0.040

0.15−0.40

...

8.50−9.50

...

...

...

A 387

2

0.21

0.55−0.80

0.035 0.040

0.15−0.40

0.50−0.80

...

0.45−0.60

...

...

5

0.15

0.30−0.60

0.040 0.030

0.50

4.00−6.00

...

0.45−0.65

...

...

7

0.15

0.30−0.60

0.030 0.030

1.00

6.00−8.00

...

0.45−0.65

...

...

9

0.15

0.30−0.60

0.030 0.030

1.00

8.00−10.0 0

...

0.90−1.10

...

...

11

0.17

0.40−0.65

0.035 0.040

0.50−0.80

1.00−1.50

...

0.45−0.65

...

...

12

0.17

0.40−0.65

0.035 0.040

0.15−0.40

0.80−1.15

...

0.45−0.60

...

...

21

0.15(b)

0.30−0.60

0.035 0.035

0.50

2.75−3.25

...

0.90−1.10

...

...

22

0.15(b)

0.30−0.60

0.035 0.035

0.50

2.00−2.50

...

0.90−1.10

...

...

A 204

A 225 A 302


Page 387


A 517

A 533

A 542

A 543

A 553


01 Sep 2005

91

0.08−0.1 2

0.30−0.60

0.020 0.010

0.20−0.50

8.00−9.50

0.40

0.85−1.05

...

V, 0.18−0.25; Nb,0.06−0.10; N, 0.03−0.07; Al, 0.04

A

0.15−0.2 1

0.80−1.10

0.035 0.040

0.40−0.80

0.50−0.80

...

0.18−0.28

...

B, 0.0025

B

0.15−0.2 1

0.70−1.00

0.035 0.040

0.20−0.35

0.40−0.65

...

0.15−0.25

...

B, 0.0005−0.005

C

0.10−0.2 0

1.10−1.50

0.035 0.040

0.15−0.30

...

...

0.20−0.30

...

B, 0.001−0.005

E

0.12−0.2 0

0.40−0.70

0.035 0.040

0.20−0.35

1.40−2.00

...

0.40−0.60

...

B,0.0015, 0.005

F

0.10−0.2 0

0.60−1.00

0.035 0.040

0.15−0.35

0.40−0.65 0.70−1.00 0.40−0.60

...

B,0.0005−0.006

H

0.12−0.2 1

0.95−1.30

0.035 0.040

0.20−0.35

0.40−0.65 0.30−0.70 0.20−0.30

...

B,0.0005

J

0.12−0.2 1

0.45−0.70

0.035 0.040

0.20−0.35

...

0.50−0.65

...

B,0.001−0.005

M

0.12−0.2 1

0.45−0.70

0.035 0.040

0.20−0.35

...

1.20−1.50 0.45−0.60

...

B, 0.001−0.005

P

0.12−0.2 1

0.45−0.70

0.035 0.040

0.20−0.35

0.85−1.20 1.20−1.50 0.45−0.60

...

B, 0.001−0.005

Q

0.14−0.2 1

0.95−1.30

0.035 0.040

0.15−0.35

1.00−1.50 1.20−1.50 0.40−0.60

...

V, 0.03−0.08

S

0.10−0.2 0

1.10−1.50

0.035 0.040

0.15−0.40

...

...

0.10−0.35

...

Ti, 0.06; Nb, 0.06

T

0.08−0.1 4

1.20−1.50

0.035 0.010 0.040−0.60

...

...

0.45−0.60

...

B, 0.001−0.005; V, 0.03−0.08

A

0.25

1.15−1.50

0.035 0.040

0.15−0.40

...

...

0.45−0.60

...

...

B

0.25

1.15−1.50

0.035 0.040

0.15−0.40

...

0.40−0.70 0.45−0.60

...

...

C

0.25

1.15−1.50

0.035 0.040

0.15−0.40

...

0.70−1.00 0.45−0.60

...

...

D

0.25

1.15−1.50

0.035 0.040

0.15−0.40

...

0.20−0.40 0.45−0.60

A

0.15

0.30−0.60

0.025 0.025

0.50 0.50

...

2.00−2.50

0.40

2.00−2.50

0.25

2.75−3.25

0.25

...

...

0.90−1.10

0.40

V, 0.03

0.90−1.10

0.25

V, 0.02

0.90−1.10

0.25

V, 0.2−0.3; Ti, 0.015−0.35; B, 0.001−0.003

B

0.11−0.1 5

0.30−0.60

0.015 0.15

C

0.10−0.1 5

0.30−0.60

0.025 0.025

0.13

B

0.23

0.40

0.035 0.040

0.20−0.40

1.50−2.00 2.60−3.25 0.45−0.60 (b)

...

V, 0.03

C

0.23

0.40

0.020 0.020

0.20−0.40

1.20−1.80 2.25−3.25 0.45−0.60 (b)

...

V, 0.03

I

0.13

0.90

0.035 0.040

0.15−0.40

...

8.50−9.50

...

...

...

II

0.13

0.90

0.035 0.040

0.15−0.40

...

7.50−8.50

...

...

...

A 645

...

0.13

0.30−0.60

0.025 0.025

0.20−0.40

...

4.75−5.25 0.20−0.35

...

Al, 0.02; N, 0.020

A 734

A

0.12

0.45−0.75

0.035 0.015

0.40

0.90−1.20 0.90−1.20 0.25−0.40

...

Al, 0.06

A 735

...

0.06

1.20−2.20(b)

0.04 0.025

0.40

A 736

A

0.07

0.40−0.70

0.025 0.025

0.40

C

0.07

1.30−1.65

0.025 0.025

0.40

...

...

0.20

0.7−1.20

0.035 0.040

0.40−0.80

0.50−1.00

A 782


...

...

0.23−0.47 0.20−0.3 Nb, 0.03−0.09 5(c)

0.60−0.90 0.70−1.00 0.15−0.25 1.00−1.3 Nb, 0.02 min 0 0.70−1.00 0.15−0.25 1.00−1.3 Nb, 0.02 min 0 ...

0.20−0.60

...

Zr, 0.04−0.12

Page 388


A 832

...


0.10−0.1 5

0.30−0.60

0.025 0.025

0.10

2.75−3.25

01 Sep 2005

...

...

0.90−1.10

V, 0.20−0.30; Ti, 0.015−0.035; B,0.001−0.003

A 844 ... 0.13 0.90 0.020 0.020 0.15−0.40 ... ... ... ... 8.50−9.50 (a) When a single value is shown, it is a maximum limit, except where specified as a minimum limit. (b) Limiting values may vary with plate thickness. (c) When specified

Mechanical tests of pressure vessel steel plate involve a minimum of one tensile test for each as-rolled plate or a minimum of two tensile tests for quenched and tempered plates. The mechanical property requirements given in ASTM specifications for pressure vessel steel plate are listed in Table 10 . Table 10 ASTM specifications of mechanical properties for pressure vessel plate made of carbon steel, HSLA steel, or low-alloy steel Tensile strength(a)

Yield strength(a) ksi

Minimum elongation(b) in 200 mm (8 in.),%

Minimum elongation(b)in 50 mm (2 in.), %

165

24

27

30

50−70

185

27

25

28

380−515

55−75

205

30

23

27

515−655

75−95

275(b)

40(b)

16

19

380−515

55−75

205

30

21

26

60

415−550

60−80

220

32

20

23

A 455

...

485−620(b)

70−90(b)

240(b)

35(b)

15

22

A 515

55

380−515

55−75

205

30

23

27

32

21

25

Material grade or type

MPa

ksi

MPa

A

310−450

45−65

B

345−485

C A 299

...

A 442

55

ASTM specification Carbon steel A 285

A 516

A 537

60

415−550

60−80

220

65

450−585

65−85

240

35

19

23

70

485−620

70−90

260

38

17

21

55

380−515

55−75

205

30

23

27

60

415−550

60−80

220

32

21

25

35

19

23

65

450−585

65−85

240

70

485−620

70−90

260

38

17

21

1

450−585(b)

65−85(b)

310(b)

45(b)

18

22

2

485−620(b)

70−90(b)

315(b)

46(b)

...

20

55(b)

...

22

2

515−655(b)

75−95(b)

380(b)

A 562

...

380−515

55−75

205

30

22

26

A 612

...

560−695(b)

81−101(b)

345

50

16

22

A 662

A

400−540

58−78

275

40

20

23

B

450−585

65−85

275

40

20

23

43

18

22

A 724 A 738

C

485−620

70−90

295

A, C

620−760

90−110

485

70

...

19

B

655−795

95−115

515

75

...

17

A

515−655

75−95

310

45

...

20

60

...

20

...

20

B

585−705

85−102

415

C

485−620

70−90

315

46

A 734

B

530−670

77−97

450

65

...

20

A 737

B

485−620

70−90

345

50

18

23

C

550−690

80−100

415

60

18

23

HSLA steel


Page 389


A 841


01 Sep 2005

...

450−585(b)

65−85(b)

310(b)

45(b)

18

22

A 202

A

515−655

75−95

310

45

16

19

B

585−760

85−110

325

47

15

18

A 203

A, D

450−585

65−85

255

37

19

23

B, E

485−620

70−90

275

40

17

21

F

515−655

75−95

345

50

...

20

A

450−585

65−85

255

37

19

23

40

17

21

Low-alloy steel

A 204

A 225

A 302

B

485−620

70−90

275

C

515−655

75−95

295

43

16

20

A

485−620

70−90

275

40

17

21

B

515−655

75−95

295

43

16

20

C

725−930

105−135

485

70

...

20

55

...

19

D

515−690

75−100

380

A

515−655

75−95

310

45

15

19

B

550−690

80−100

345

50

15

18

C, D

550−690

80−100

345

50

17

20

75

...

20

A 353

...

690−825

100−120

515

A 387

2, 12 (class 1)

380−550

55−80

230

33

18

22

11 (class 1)

415−585

60−85

240

35

19

22

22, 21, 5, 7, and 9 (class 1)

415−585

60−85

205

30

...

18

2 (class 2)

485−620

70−90

310

45

18

22

11 (class 2)

515−690

75−100

310

45

18

22

40

19

22

A 517

A 533

A 542

12 (class 2)

450−585

65−85

275

22, 21, 5, 7, and 9 (class 2)

515−690

75−100

310

45

...

18

91

585−760

85−110

415

60

...

18

A, B, C, F, H, J, M, S, T

795−930

115−135

690

100

...

16

E, P, Q

725−930(b)

105−135(b)

620(b)

90(b)

...

14

1

550−690

80−100

345

50

...

18

2

620−795

90−115

485

70

...

16

3

690−860

100−125

570

83

...

16

85

...

14

1

725−860

105−125

585

2

795−930

115−135

690

100

...

13

3

655−795

95−115

515

75

...

20

4

585−760

85−110

380

55

...

20

60

...

18

4a

585−760

85−110

415

1

725−860

105−125

585

85

...

14

2

795−930

115−135

690

100

...

14

3

620−795

90−115

485

70

...

16

A 553

I, II

690−825

100−125

585

85

...

20

A 645

...

655−795

95−115

450

65

...

20

A 734

A

530−670

77−97

450

65

...

20

A 735

1(c)

550−690

80−100

450

65

12

18

2(d)

585−720

85−105

485

70

12

18

A 543


Page 390


A 736

A 782

A 832


01 Sep 2005

3

620−750

90−110

515

75

12

18

4

655−790

95−115

550

80

12

18

A1

620−760

90−110

550

80

...

20

A2

415−550(b)

60−80(b)

345(b)

50(b)

...

20

A3

485−620(b)

70−90(b)

415(b)

60(b)

...

20

90

...

20

C1

690−825

100−120

620

C3

620−760(b)

90−110

550(b)

80(b)

...

20

1

670−820

97−119

550

80

...

18

2

740−890

107−129

620

90

...

17

100

...

16

60

...

18

3

795−940

115−136

690

...

585−760

85−110

415

A 844 ... 585 85 ... 20 690−825 100−120 (a) Where a single value is shown, it is a minimum. (b) Minimum and/or maximum values depend on plate thickness. (c) As-rolled class 1 plate is limited to 25 mm (1 in.) thickness. (d) As-rolled and aged class 2 plate is limited to 25 mm (1 in.) thickness.

Aircraft quality plates are used for important or highly stressed parts of aircraft, missiles, and other applications involving stringent requirements. Plates of this quality require exacting steelmaking, conditioning, and process controls and are generally furnished from electric furnace steels in order to meet the internal cleanliness requirements outlined in Aerospace Materials Specifications AMS-2301. The primary requirements of this quality are a high degree of internal soundness, good uniformity of chemical composition, good degree of cleanliness, and a fine austenitic grain size. Aircraft quality plates can be supplied in the hot-rolled or thermally treated condition. Forging quality plates are intended for forging, quenching and tempering, or similar purposes or when uniformity of composition and freedom from injurious imperfections are important (see ASTM A 827). Plates of this quality are produced from killed steel and are ordinarily furnished with the phosphorus content limited to 0.035% maximum and the sulfur content limited to 0.040% maximum by heat analysis. Table 11 lists some AISI/SAE steels suitable for forging quality plate. Plates of this quality can be produced to chemical ranges and limits and mechanical properties. When mechanical properties are specified, two tension tests from each heat are taken from the same locations at tests for structural quality. Factors affecting mechanical properties are discussed in the section "Mechanical Properties" in this article. Table 11 Compositions of forging quality steel plate specified in ASTM A 827 Grade

Element, %

UNS

AISI

C

Mn

P(a)

S(a)

Si

G10090

1009

0.15(a)

0.60(a)

0.035

0.040

0.15−0.40

G10200

1020

0.17−0.23

0.30−0.60

0.035

0.040

0.15−0.40

G10350

1035

0.31−0.38

0.60−0.90

0.035

0.040

0.15−0.40

G10400

1040

0.36−0.44

0.60−0.90

0.035

0.040

0.15−0.40

G10450

1045

0.42−0.50

0.60−0.90

0.035

0.040

0.15−0.40

1050

0.47−0.55

0.60−0.90

0.035

0.040

0.15−0.40

G10500 (a) Maximum

Mechanical Properties Of the various mechanical properties normally determined for steel plate, yield strength is an important design criterion in structural applications. Tensile strength is also an important design consideration in many design codes in the United States, but is useful primarily as an indication of fatigue properties. Yield strength is a design criterion in most design codes when the ratio of yield to tensile strength is less than 0.5. Ductility, as measured by tensile elongation and reduction in area, is seldom in itself a valuable design criterion, but is sometimes used as an indication of toughness and suitability for certain applications. The mechanical properties of steel plate in the hot-finished condition are influenced by several variables, of which chemical composition is the most influential. Other factors include deoxidation practice, finishing temperature, plate thickness, and the presence of residual elements such as nickel, chromium, and molybdenum. For steels used in the hot-finished condition (such as plate), carbon content is the single most important factor in determining mechanical properties. The static tensile properties of the various grades, types, and classes of steel plate covered by ASTM specifications are listed in Tables 7 and 10 . It should be noted that some of these values vary with plate thickness and/or width. An example of the variation of tensile strength and elongation with thickness is shown in Fig. 1 , which presents the minimum expected values for 0.20% C steel plate from 13 to 125 mm (1=2to 5 in.) thick. Plate under 13 mm (1=2in.) thick would show even slightly higher tensile strength and lower elongation because of the increased amount of hot working during rolling and the faster cooling rates


Page 391



01 Sep 2005

after rolling. Fig. 1 Effect of thickness on tensile properties of 0.20% C steel plate

The distribution of the tensile properties obtained for a larger number of heats of A 285, A 515, and A 516 steel plate is illustrated in Fig. 2 , which also shows the distribution of the carbon and manganese content. The use of the carbon and manganese contents to control mechanical properties is clearly shown in Fig. 2 ; higher carbon and manganese contents accompany higher yield strengths. Fig. 2 Distribution of tensile properties and chemical composition of carbon steel plate. Data represent all the as-hot-rolled plate, 6 to 50 mm (1=4to 2 in.) thick, purchased to these specifications by one fabricator during a period of 8 years.


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01 Sep 2005

Figure 3 repeats the data in Fig. 2 for the 224 heats of A 285, grade C plate. However, the data are augmented by the individual distributions for the various ranges of plate thickness included in the investigation. When steel is produced to a mechanical property requirement, it is common practice to vary the carbon and manganese to compensate for size influence. The use of higher-than-average carbon (and manganese) content to maintain yield strength as plate thickness increases is evident in Fig. 3 . Fig. 3 Distribution of tensile properties and chemical composition of ASTM A 285, grade C, carbon steel plate. Data represent all the as-hot-rolled plate (224 heats from 6 mills) purchased to this specification by one fabricator during a period of 8 years.


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01 Sep 2005

The common mechanical properties of hot-finished steel, including plate, reliably related to each other, and this relation is relatively free from influence of composition for most of these properties. Figure 4 shows the relationship between yield strength, elongation, and tensile strength over a wide range of tensile strengths for various hot-rolled carbon steels. Fig. 4 Relation of tensile properties for hot-rolled carbon steel


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01 Sep 2005

Residual alloying elements generally have a minor strengthening effect on hot-finished steels, such as plate. This effect cannot be considered in design because residuals vary greatly among the different steel producing plants. This influence is shown in Fig. 5 , which demonstrates that the effect is minor. Fig. 5 Effect of carbon and amount of residuals on tensile properties of hot-finished carbon steel. Curves marked high residuals represent steel containing 0.06 to 0.12% Ni, 0.06 to 0.13% Cr, and 0.08 to 0.13% Mo. Curves marked low residuals represent steel containing 0.05% Ni max, 0.05% Cr max, and 0.04% Mo max. Total of 58 heats tested


Page 395



01 Sep 2005

Hardness is a relatively simply test to perform and is closely related to tensile strength, as shown in Fig. 6 . A simple hardness test, used in conjunction with the data in Fig. 4 , can be used to estimate yield strength, elongation, and tensile strength. Fig. 6 Relation between hardness and tensile strength of steel. Range up to 300 HB is applicable to the hot-finished steel discussed in this article.

Fatigue Strength. The high-cycle (>1 million) fatigue properties of hot-finished steel, often called the fatigue limit, are more or less directly related to tensile strength and are greatly affected by the surface condition. The fatigue limit of machined specimens is about 40% of the tensile strength, depending on the surface finish. In contrast, unmachined hot-rolled steel, when loaded so that fatigue stresses are concentrated at the surface, will have a considerably lower fatigue limit because of decarburization, surface roughness, and other surface imperfections. For this reason, the location of maximum fatigue stresses should be carefully considered; for structural members designed in hot-finished steel, the surface should be machined off from critically stressed areas or an allowance made for the weakness of the hot-finished surface. The presence of inclusions in hot-finished steel may also have an adverse effect on the fatigue limit. Large inclusions are considered harmful under the dynamic stresses of impact or fatigue, and the effect is greater in the harder steels. Low-Temperature Impact Energy. When notch toughness is an important consideration, satisfactory service performance can be ensured by proper selection of the steel that will behave in a tough manner at its lower operating temperature. The Charpy V-notch tests and crack-starter drop-weight tests provide a fairly reliable indication of the tendency toward brittle fracture in service. The transition temperatures of hot-finished steels are controlled principally by their chemical composition and ferrite grain size. For the steels considered in this article, carbon is of primary importance because of its effect is raising the transition temperature, lowering the maximum energy values, and widening the temperature range between completely tough and completely brittle behavior.


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01 Sep 2005

Manganese (up to about 1.5%) improves low-temperature properties. Also, as mentioned previously, the transition temperature is affected by the deoxidation practice used. The transition temperature decreases and the energy absorption before fracture at normal temperatures increases in the order of rimmed, capped, semikilled, and killed steels. In addition, killed steels contain larger amounts of silicon or aluminum than semikilled steels, and these elements improve low-temperature toughness and ductility. Because of variations in finishing temperatures and cooling rates, plate thickness influences the grain size and therefore the transition temperature. Extensive data on the impact properties of hot-finished steel are given in the article "Notch Toughness of Steels" in this Volume. Elevated-Temperature Properties. The steel plate used in pressure vessel applications is often subjected to long-term elevated temperatures. Of the carbon and low-alloy steels used for pressure vessel plate, the behavior of 2.25Cr-1Mo steel (ASTM A 387, Class 22, in Table 9 ) at elevated temperatures has been studied more thoroughly than any other steel and has become the reference for comparing the elevated-temperature properties of low-alloy steels. Further information on the elevated-temperature properties of 2.25Cr-1Mo steel can be found in the article "Elevated-Temperature Properties of Ferritic Steels" in this Volume. Directional Properties. An important characteristics of steel plate, known as directionality or fibering, must be considered. During the rolling operations, many inclusions, which are in a plastic condition at rolling temperatures, are elongated in the direction of rolling. At the same time, localized chemical segregates that have formed during solidification of the steel are also elongated. These effects reduce the ductility and impact properties transverse to the rolling direction, but have little or no effect on strength.

Fabrication Considerations Formability. The cold formability of steel plate is directly related to the yield strength and ductility of the material. The lower the yield strength, the smaller the load required to produce permanent deformation; high ductility allows large deformation without fracture. Therefore, the lower-carbon grades are most easily formed. Operations such as shearing and blanking are usually limited by the lack of the available facilities as the plate thickness increases. This also applies to bending operations. Of course, an adequate bend radius must be used to avoid fracture. Because of fibering effects, the direction of bend is also important; when the axis of a bend is parallel to the direction of rolling, small bend radii are usually difficult to form because of the danger of cracking. Machinability. Machining operations are usually performed with little difficulty on most plate steels up to about 0.50% C. Higher-carbon steels can be annealed for softening. Steels with low carbon and manganese content, such as 1015, with large quantities of free ferrite in the microstructure may be too soft and gummy for good machining. Increasing the carbon content (to a steel such as 1025) improves the machinability. Machining characteristics can be improved by factors that break up the chip as it is removed. This is usually accomplished by the introduction of large numbers of inclusions such as manganese sulfides or complex oxysulfides. These "free-machining" steels are somewhat more expensive, but are cost-effective when extensive machining is involved. Weldability is a relative term that describes the ease with which sound welds possessing good mechanical properties can be produced in a material. The chief weldability factors are composition, heat input, and rate of cooling. These factors produce various effects, such as grain growth, phase changes, expansion, and contraction, which in turn determine weldability. Heat input and cooling rate are characteristics of the specific process and technique used and the thickness of the metal part being welded. Therefore, weldability ratings should state the conditions under which the rating was determined and the properties and soundness obtained. For carbon steels, the carbon and manganese contents are the primary elements of the composition factor that determine the effect of the steel of given heating and cooling conditions. The great tonnage of steel used for welded applications consists of low-carbon steel, up to 0.30% C. Generally, steels with a carbon content less than 0.15% are readily weldable by any method. Steel with a carbon range of 0.15 to 0.30% can usually be welded satisfactorily without preheating, postheating, or special electrodes. For rather thick sections (>25 mm, or 1 in.), however, special precautions such as 40 °C (100 °F) minimum preheat, 40 °C (100 °F) minimum temperature between weld passes, and a 540 to 675 °C (1000 to 1250 °F) stress relief may be necessary. Higher-carbon and higher-manganese grades can often be welded satisfactorily if preheating, special welding techniques, and postheating and peening are used. In the absence of such precautions to control the rate of cooling and to eliminate high stress gradients, cracks may occur in the weld and base metal. In addition, base metal properties such as strength, ductility, and toughness may be greatly reduced. All comments about the effect of carbon and manganese on weldability must be qualified in terms of section size because of its relationship to heat input and cooling rate. In welding thicker sections, such as plate, the relatively cold base metal serves to greatly accelerate the cooling rate after welding with the result that plate thickness is a very important consideration. Figure 7 shows the effect of plate thickness and carbon equivalent on weldability as expressed in terms of a notch bend test. Fig. 7 Ratio (welded to unwelded) of bend angle for normalized steel plate. A high value of the ratio indicates high weldability. Source: Ref 2


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01 Sep 2005

REFERENCES 1. Plates; Rolled Floor Plates: Carbon, High Strength Low Alloy, and Alloy Steel, AISI Steel Products Manual, American Iron and Steel Institute, 1985 2. Weldability of Steels, Welding Research Council, 1953


Page 398


Hot-Rolled Steel Bars and Shapes

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Carbon and Low-Alloy Steels Hot-Rolled Steel Bars and Shapes Revised by Timothy E. Moss, J.M. Hambright, and T.E. Murphy, Inland Bar and Structural, Division of Inland Steel Company;and J.A. Schmidt, Joseph T. Ryerson and Sons, Inc. HOT-ROLLED STEEL BARS and other hot-rolled steel shapes are produced from ingots, blooms, or billets converted from ingots or from strand cast blooms or billets and comprise a variety of sizes and cross sections. Bars and shapes are most often produced in straight lengths, but bars in some cross sections in smaller sizes are also produced in coils. The term "bar" includes: • Rounds, squares, hexagons, and similar cross sections 9.5 mm (3=8in.) and greater across • Flats greater than 5.16 mm (0.203 in.) in thickness and 152 mm (6 in.) and less in width, or 5.84 mm (0.230 in.) and greater in thickness and 203 mm (8 in.) and less in width • Small angles, channels, tees, and other standard shapes less than 76 mm (3 in.) across • Concrete-reinforcing bars The term "shape" includes structural shapes and special shapes. Structural shapes are flanged, are 76 mm (3 in.) or greater in at least one cross-sectional dimension, and are used in structures such as bridges, buildings, ships, and railroad cars. Special shapes are those designed by users for specific applications.

Dimensions and Tolerances The nominal dimensions of hot-rolled steel bars and shapes are designated in inches or millimeters with applicable tolerances, as shown in ASTM A 6 and A 29. Bars with certain quality descriptors have size limitations; these are covered in discussions of individual product qualities later in this article. Bars or shapes can be cut to length in the mill by a number of methods, such as hot or cold shearing or sawing. The method used is determined by cross section, grade, and customer requirements. Some end distortion is to be expected from most methods. When greater accuracy in length or freedom from distortion is required, bars of shapes can be cut overlength, then recut on one or both ends by cold sawing or equivalent means. If a bar or shape requires straightening, prior annealing is sometimes necessary, depending on the grade of steel and the cross-sectional shape of the part. The processing necessary to meet straightness tolerances is not intended to improve either the surface finish or accuracy of cross-sectional shape and may result in increased surface hardness. Length and straightness tolerances for bars and shapes are found in ASTM A 6 and A 29.

Surface Imperfections Most carbon steel and alloy steel hot-rolled bars and shapes contain surface imperfections with varying degrees of severity. In virtually all cases, these defects are undesirable and may in some applications affect the integrity of the finished product. Included in the manufacturing process for hot-rolled bars and shapes are various steps designed to minimize or eliminate surface defects. These steps include inspection of both the semifinished and the finished product and either subsequent removal of the defects or rejection of the material if defect removal is not possible. Inspection techniques range from visual inspection of the semifinished material to sophisticated electronic inspection of the finished product. Defects found in the semifinished product can be removed by hot scarfing, grinding, or chipping. Defects in the finished products are generally removed by grinding, turning, or peeling and, to a lesser degree, by chipping. Currently, it is not technically feasible to produce defect-free hot-rolled bars. With the current demand for high-quality bar products, it is becoming increasingly common to subject hot-rolled bars to a cold-finishing operation, such as turning or grinding, coupled with a sensitive electronic inspection. With this process route, it becomes possible to significantly reduce both the frequency and the severity of surface defects. Seams, Laps, and Slivers


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01 Sep 2005

Seams, laps, and slivers are probably the most common defects in hot-rolled bars and shapes. Seams are longitudinal defects that can vary greatly in length and depth. It is quite common for steel users to refer to any longitudinal defect as a seam regardless of the true nature of the defect. However, there is a classical definition of a seam, as follows. Gas comes out of the solution as the liquid steel solidifies. This gas is trapped as bubbles or blowholes by the solidifying steel and appears as small holes under the surface of the steel. When the steel is reheated, some areas of the surface may scale off, exposing and oxidizing the interior of these blowholes. This oxidation prevents the blowholes from welding shut during rolling. This rolling then elongates the steel, resulting in a longitudinal surface discontinuity⎯a seam. As viewed in the cross section, seams are generally characterized as being perpendicular to the surface, completely surrounded by decarburization, and associated with disperse oxides. Laps are mechanical defects that occur during the hot rolling of both semifinished and finished material. Laps are nothing more than a folding over of the material, resulting, for example, from gouging during the rolling process or misalignment of the pass lines or rolls. As viewed in the cross section, laps are characterized as being at an angle from the steel surface; they have decarburization on one side only of the defect and often contain entrapped scale. Slivers usually appears as a scablike defect, adhering on one end to the surface of the hot-rolled steel. They are normally pressed into the surface during hot rolling. They can originate from short, rolled out defects such as torn corners that are not removed during conditioning. They can also result from conditioning gouges or mechanical gouges during rolling. Although there is no specific metallographic definition of slivers, metallographic examination can be used to determine the origin of these defects. Decarburization Another condition that could be considered a surface defect is decarburization. This condition is present to some degree on all hot-rolled steel. Decarburization occurs at very high temperatures when the surface carbon of the steel reacts with the oxygen in the furnace atmosphere. This loss of surface carbon results in a surface that is softer and unsuitable for any application involving wear or fatigue. Because of the existence of this condition, steel ordered for critical applications can be produced oversize and then ground to desired size. Allowance for Surface Imperfections in Machining Applications Experience has shown that when purchasers order hot-rolled or heat-treated bars that are to be machined, it is advisable for the purchaser to make adequate allowances for the removal of surface imperfections and to specify the sizes accordingly. These allowances depend on the way the surface metal is removed, the length and size of the bars, the straightness, the size tolerance, and the out-of-round tolerance. Bars are generally straightened before machining. For special quality carbon steel bars and regular quality alloy steel bars, either resulfurized or nonresulfurized (see the article "Cold-Finished Steel Bars" in this Volume), it is advisable that allowances for centerless-turned or centerless-ground bars be adequate to permit stock removal of not less than the amount shown below: Recommended minimum machining allowance per side, % of specific size

Bar diameter mm

in.

Nonresulfurized

Resulfurized

≤51

≤2

2.6

3.4

>2

1.6

2.4

>51 Source: Ref 1

Note that these allowances are based on bars within straightness tolerance. Also, because straightness is a function of length, additional machining allowance may be required for turning long bars on centers. For steel bars subject to magnetic particle inspection, additional stock removal is recommended, as indicated in Table 1 . Table 1 Recommended minimum stock removal for steel bars subject to magnetic particle inspection Minimum stock removal from the surface(a)

Hot-rolled size mm

in.

mm

in.

0.76

0.030

Up to 12.7

Up to 1=2

>12.7−19

>1=2−3=4

1.14

0.045

>19−25

> =4−1

1.52

0.060

>25−38

>1−11=2

1.90

0.075

>38−51

>1 =2−2

2.29

0.090

3

1

1

>51−64

>2−2 =2

3.18

0.125

>64−89

>2 =2−3 =2

3.96

0.156

1

1


Page 400



01 Sep 2005

>89−114

>31=2−41=2

4.75

0.187

>114−152

>4 =2−6

6.35

0.250

7.92

0.312

9.52

0.375

11.10

0.437

1

1

>152−191

>6−7 =2

>191−229

>7 =2−9 1

>229−254 >9−10 (a) The minimum reduction in diameter of rounds is twice the minimum stock removal from the surface.

The allowances described above are usually more than sufficient to remove surface imperfections and result in considerable loss of material. Therefore, most experienced fabricators remove considerably less stock than recommended and take their chances on occasional difficulties. In conventional practice, depth of machining for hot-rolled bars is 1.6 mm (1=16in.) for bars 38 to 76 mm (11=2to 3 in.) in diameter, and 3.2 mm (1=8in.) for bars over 76 mm (3 in.) in diameter.

Surface Treatment It is uncommon for hot-rolled steel bars and shapes to be descaled by the producer or protected from the weather during transit. Most cleaning and coating operations are done either by the customer or by an intermediate processor. Descaling of hot-rolled bars and shapes is generally done by either pickling or blasting, depending on the end use. There are several subsequent coatings that can be used. Oil is both the simplest and the least expensive to use and acts as a temporary rust preventive. Lime, in addition to serving as a rust preventive, can serve as a carrier for lubricants used during cold drawing or cold forging. A more sophisticated system includes descaling, followed by a zinc phosphate coating, coupled with a dry lubricant. This system provides some rust protection and serves as a lubricant for cold-forming operations.

Heat Treatment Hot-rolled low-carbon and medium-carbon steel bars and shapes are often used in the as-rolled condition, but hot-rolled bars of higher-carbon steel and most hot-rolled alloy steel bars must be heat treated in order to attain the hardness and microstructure best suited for the final product or to make them suitable for processing. Such heat treatment consists of one or more of the following: some form of annealing, stress relieving, normalizing, quenching, and tempering. Ordinary annealing is the term generally applied to heat treatment used to soften steel. The steel is heated to a suitable temperature, held there for some period of time, and then cooled; specific times, temperatures, and cooling rates vary. Maximum hardness compatible with common practice can be specified. Annealing for specified microstructures can be performed to obtain improved machinability or cold-forming characteristics. The structures produced may consist of lamellar pearlite or spheroidized carbides. Special control of the time and temperature cycles is necessary. A compatible maximum hardness can be specified. Stress relieving involves heating to a sub-critical temperature and then cooling. For hot-rolled bars, the principal reason for stress relieving is to minimize distortion in subsequent machining. It is used to relieve the stresses resulting from cold-working operations, such as special machine straightening. Normalizing involves heating to a temperature above the critical temperature range and then cooling in air. A compatible maximum hardness can be specified. Hardening by quenching consists of heating steel to the correct austenitizing temperature, holding at that temperature for a sufficient time to produce homogeneous austenite, and quenching in a suitable medium (water, oil, synthetic oil or polymer, molten salts, or low-melting metals) depending on chemical composition and section thickness. Tempering is an operation performed on normalized or quenched steel bars. In this technique, the bars are reheated to a predetermined temperature below the critical range and then cooled under suitable conditions. When a hardness requirement is specified for normalized and tempered bars, the bars are ordinarily produced to a range of hardnesses equivalent to a 0.4 mm range of Brinell impression diameters. Quenched and tempered bars are ordinarily produced to a 0.3 mm range of Brinell impression diameters. Quenched and tempered bars can also be produced to minimum mechanical property requirements.

Product Requirements Hot-rolled steel bars and shapes can be produced to chemical composition ranges or limits, mechanical property requirements, or both. The mechanical testing of hot-rolled steel bars and shapes can include tensile, Brinell or Rockwell hardness, bend, Charpy impact, fracture toughness, and short-time elevated-temperature tests, as well as test for elastic limit, proportional limit, and offset yield strength, which require the use of an extensometer or plotting of a stress-strain curve. These tests are covered by ASTM A 370 and other ASTM standards. Other tests sometimes required include the measurement of grain size and hardenability. Austenitic grain size is determined by the McQuaid-Ehn test, which is described in ASTM E 112. This test involves metallographic examination of a carburized specimen to observe prior austenitic grain boundaries. Hardenability can be measured by several methods, the most common beingthe Jominy end-quench test, as described in ASTM A 255 (see the article "Hardenability of Carbon and Low-Alloy Steels" in this Volume).


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Soundness and homogeneity can be evaluated by fracturing. The fracture test is commonly applied only to high-carbon bearing quality steel. Location of samples, number of tests, details of testing technique, and acceptance limits based on the test should be established in each instance. Testing for nonmetallic inclusions consists of careful microscopic examination (at 100×) of prepared and polished specimens. The specimens should be taken on a longitudinal plane midway between the center and surface of the product. Location of specimens, number of tests, and interpretation of results should be established in each instance. Typical testing procedures are described in ASTM E 45. Nonmetallic inclusion content can also be measured on the macroscopic scale by magnetic particle tests such as those described in AMS 2300 and 2301. These tests involve the measurement of inclusion frequency and severity in a sampling scheme that represents the interior of the material. Surface and subsurface nonuniformities are revealed by magnetic particle testing. This test was developed for, and is used on, fully machined or ground surfaces of finished parts. When the magnetic particle test is to be applied to bar stock, short-length samples should be heat treated and completely machined or ground. Tensile and hardness tests are the most common mechanical tests performed on hot-rolled steel bars and shapes. Hardness is a relatively simple property to measure, and it is closely related to tensile strength, as shown in Fig. 1 . When Fig. 2 is used together with Fig. 1 , a simple hardness test can give an estimate of yield strength and elongation, as well as tensile strength. Fig. 1 Relationship between hardness and tensile strength of steel. Range up to 300 HB is applicable to the hot-finished steel discussed in this article. Source: Ref. 2

Fig. 2 Relation of tensile properties for hot-rolled carbon steel


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It is not practicable to set definite limitations on tensile strength or hardness for carbon or alloy steel bars in the as-rolled condition. For mill-annealed steel bars, there is a maximum tensile strength or a maximum hardness (Table 2 ) that can be expected for each grade of steel. For steel bars in the normalized condition, maximum hardness, maximum tensile strength, minimum hardness, or minimum tensile strength can be specified. For normalized and tempered bars and for quenched and tempered bars, either maximum and minimum hardness or maximum and minimum tensile strength can be specified; for either property, the range that can be specified varies with tensile strength and is equivalent to a 0.4 mm range of Brinell indentation diameters at any specified location for normalized and tempered bars and to a 0.3 mm range for quenched and tempered bars. Table 2 Lowest maximum hardness that can be expected for hot-rolled steel bars, billets, and slabs with ordinary mill annealing Maximum hardness, HB(a) Steel grade

Straightened

Nonstraightened

1141

201

192

1144

207

197

1151

207

201

1541

207

197

1548

212

207

1552

212

207

15B41

207

197

15B48

212

207

1330

187

179

1335

197

187

Carbon steels

Alloy steels


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1340

201

192

1345

212

201

4012

149

143

4023

156

149

4024

156

149

4027

170

163

4028

170

163

4037

192

183

4047

212

201

4118

170

163

4130

183

174

4137

201

192

4140

207

197

4142

212

201

4145

217

207

4147

223

212

4150

235

223

4161

241

229

4320

207

197

4340

235

223

4419

170

163

4615

174

167

4620

179

170

4621

179

170

4626

187

179

4718

179

170

4720

170

163

4815

223

192

4817

229

197

4820

229

197

5015

156

149

50B44

207

197

50B46

217

201

50B50

217

201

50B60

229

217

5120

170

163

5130

183

174

5132

187

179

5135

192

183

5140

197

187

5145

229

197

5147

217

207

5150

212

201

5155

229

217

5160

235

223

51B60

235

223

6118

163

156

6150

217

207

81B45

201

192


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01 Sep 2005

8615

163

156

8617

163

156

8620

170

163

8622

179

170

8625

179

170

8627

183

174

8630

187

179

8637

201

192

8640

207

197

8642

212

201

8645

217

207

8655

235

223

8720

170

163

8740

212

201

8822

187

179

9254

241

229

9255

241

229

9260

248

235

94B17

156

149

94B30 183 (a) Specific microstructure requirements may necessitate modification of these hardness numbers.

174

It is essential that the purchaser specify the positions at which hardness readings are to be taken. When both hardness and tensile strength are specified at the same position, the limits should be consistent with each other. When hardness limits are specified as surface values, they may be inconsistent with tensile-test values, which of necessity are properties of the bulk metal; the inconsistency will vary according to the size of the bar and the hardenability of the steel. The purchaser should specify limits that take this inconsistency into account. If the locations of hardness readings are not specified on the purchaser's order or specification, the hardness values are applicable to the bar surface after removal of decarburization. Hardness correction factors for bars of various diameters as described in ASTM E 18 should be employed if a flat area is not available on the bar tested. Generally, yield strength, elongation, and reduction in area are specified as minimums for steel only in the quenched and tempered or the normalized and tempered condition, and they should be consistent with ultimate tensile strength or hardness. When quenched and tempered bars are cold worked by cold straightening, stress relieving may be required to restore elastic properties and to improve ductility.

Product Categories Hot-rolled carbon steel bars are produced to two primary quality levels: merchant quality and special quality. Merchant quality is the lower quality level and is not suitable for any operation in which internal soundness or freedom from surface imperfections is of primary importance. Special, quality includes all bar categories with end-use-related and restrictive quality requirements. The mechanical properties of hot-rolled carbon steel bars in the as-rolled condition are influenced by: • Chemical composition • Thickness or cross-sectional area • Variables in mill design and mill practice Carbon content is the dominant factor. The minimum expected mechanical properties of commonly used grades of hot-rolled carbon steel bars are shown in Fig. 3 . Fig. 3 Estimated minimum tensile properties of selected hot-rolled carbon steel bars


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Quality descriptors for hot-rolled alloy steel bars are related to suitability for specific applications. Characteristics considered include inclusion content, uniformity of chemical composition, and freedom from surface imperfections. Carbon steel and alloy steel structural shapes and special shapes do not have specific quality descriptors but are covered by several ASTM specifications (Table 3 ). In most cases, these same specifications also cover structural quality steel bars. The ASTM specifications covering other qualities of hot-rolled bars are listed in Table 4 . The various categories of hot-rolled steel


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bar products and their characteristics are described in the following sections. Table 3 Typical ASTM specifications for structural quality steel bars and steel structural shapes Covered in ASTM A 6 Specification


Carbon steels A 36(a)(b)

Carbon steel plates, bars, and shapes

A 131(c)

Carbon and HSLA steel plates, bars, shapes, and rivets for ships

A 529

Carbon steel plates, bars, shapes, and sheet piling with minimum yield strength of 290 MPa (42 ksi)

A 709

Carbon, alloy, and HSLA steel plates, bars, and shapes for bridges

Alloy steel A 710

Age-hardening low-carbon Ni-Cu-Cr-Mo-Nb and Ni-Cu-Nb alloy steel plates, bars, and shapes

High-strength low-alloy (HSLA) steels A 131(c)

See above under Carbon Steel

A 242

HSLA steel plates, bars, and shapes

A 572

Nb-V HSLA steel plates, bars, shapes, and sheet piling

A 588

HSLA steel plates, bars, and shapes with minimum yield point of 345 MPa (50 Ksi)

A 633

Normalized HSLA steel plates, bars, and shapes

A 690

HSLA steel H-piles and sheet piling for use in marine environments (a) This ASTM specification is also published by the American Society of Mechanical Engineers, which adds an S in front of the A. (b) See also Canadian Standards Association (CSA) specification G40.8. (c) See also Section 39 of the ABS specifications.

Table 4 Typical ASTM specifications for hot-rolled steel bars See Table 3 for ASTM specifications for structural quality bars and structural shapes. Specification


Carbon steels A 321(a)

Quenched and tempered carbon steel bars

A 575(a)

Merchant quality carbon steel bars

A 576(a)

Special quality carbon steel bars

A 663(a)

Merchant quality carbon steel bars subject to mechanical property requirements

A 675(a)

Special quality carbon steel bars subject to mechanical property requirements

Alloy steels A 295

Bearing quality high-carbon chromium steel billets, forgings, tube rounds, bars, rods, and tubes

A 304(a)

Alloy steel bars subject to end-quench hardenability requirements

A 322(a)

Alloy steel bars for regular constructional applications

A 434(a)

Quenched and tempered alloy steel bars, hot rolled or cold finished

A 485

Bearing quality high-carbon chromium steel billets, tube rounds, bars, and tubes modified for high hardenability

A 534

Carburizing alloy steel billets, tube rounds, bars, rods, wire, and tubes of bearing quality

A 535

Special quality alloy steel billets, bars, tube rounds, rods, and tubes for the manufacture of antifriction bearings

(a) Covered in ASTM A 29

Merchant Quality Bars Merchant quality is the least restrictive descriptor for hot-rolled carbon steel bars. Merchant quality bars are used in the production of noncritical parts of bridges, buildings, ships, agricultural implements, road-building equipment, railway equipment,


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and general machinery. These applications require only mild cold bending, mild hot forming, punching, and welding. Mild cold bending is bending in which a generous bend radius is used and in which the axis of the bend is at right angles to the direction of rolling. Merchant quality bars should be free from visible pipe; however, they may contain pronounced chemical segregation, and for this reason, product analysis tolerances are not appropriate. Internal porosity, surface seams, and other surface irregularities may be present and are generally expected in bars of this quality. Consequently, merchant quality bars are not suitable for applications that involve forging, heat treating, or other operations in which internal soundness or freedom from surface imperfections is of primary importance. Grades. Merchant quality bars can be produced to meet both chemical composition (heat analysis only) and mechanical properties. These steels can be supplied to chemical compositions within the ranges of 0.50% C (max), 0.60% Mn (max), 0.04% P (max), and 0.05% S (max), but are not produced to meet any specific silicon content, grain size, or any other requirement that would dictate the type of steel produced. Merchant quality steel bars do not require the chemical ranges typical of standard steels. They are produced to wider carbon and manganese ranges and are designated by the prefix "M." When ordering merchant quality bars to meet mechanical properties, the following strength ranges are to be used up to a maximum of 655 MPa (95 ksi): • 70 MPa (10 ksi) for minimums up to but not including 415 MPa (60 ksi) • 80 MPa (12 ksi) for minimums from 415 MPa (60 ksi) up to but not including 460 MPa (67 ksi) • 100 MPa (15 ksi) for minimums from 460 to 550 MPa (67 to 80 ksi) Specification ASTM A 663 defines the requirements for hot-wrought merchant quality carbon steel bars and bar-size shapes intended for noncritical constructional applications. Sizes. Merchant quality steel rounds are not produced in diameters greater than 76 mm (3 in.).

Special Quality Bars Special quality bars are employed when end use, method of fabrication, or subsequent processing treatment requires characteristics not available in merchant quality bars. Typical applications, including many structural uses, require hot forging, heat treating, cold drawing, cold forming, and machining. Special quality bars are required to be free from visible pipe and excessive chemical segregation. Also, they are rolled from billets that have been inspected and conditioned, as necessary, to minimize surface imperfections. Frequency and degree of surface imperfections are influenced by chemical composition, type of steel, and bar size. Resulfurized grades, certain low-carbon killed steels, and boron-treated steels are most susceptible to surface imperfections. Some end uses or fabricating procedures can necessitate one or more extra requirements. These requirements include special hardenability, internal soundness, nonmetallic inclusion rating, and surface condition and are described in the AISI manual covering hot-rolled bars. The quality descriptorfor bars to which only one of these special requirements is applied is Restrictive Requirement Quality A. When a single special restriction other than the four mentioned above is applied, the quality descriptor is Restrictive Requirement Quality B. Multiple Restrictive Requirement Quality bars are those to which two or more restrictive requirements are applied. Special quality steel bars can be produced using rimmed, capped, semikilled, or killed deoxidation practice. The appropriate type is dependent on chemical composition, quality, and customer specifications. Killed steels can be produced to coarse or fine austenitic grain size. Special quality steel bars are produced to product chemical composition tolerances and can be purchased on the basis of heat composition. Special quality steel bars can also be produced to meet mechanical property requirements. The tensile strength ranges are identical to those presented in the section "Merchant Quality Bars" in this article. Additional information on mechanical property requirements and test frequencies is available in the appropriate ASTM specifications. Sizes. Special quality steel bars are commonly produced in the following sizes: • • • • •

Rounds: 6.4 to 254 mm (1=4to 10 in.) Squares: 6.4 to 154 mm (1=4to 61=16in.) Round-cornered squares: 9.5 to 203 mm (3=8to 8 in.) Hexagons: 9.5 to 103 mm (3=8to 41=16in.) Flats: greater than 5.16 mm (0.203 in.) in thickness and 152 mm (6 in.) and less in width, or 5.84 mm (0.230 in.) and greater in thickness and 203 mm (8 in.) and less in width

Common size ranges have not been established for special quality bars of other shapes, including bar-size shapes, ovals, half-ovals, half-rounds, octagons, and special bar-size shapes.

Carbon Steel Bars for Specific Applications


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Cold-working quality is the descriptor (replacing the older terminology of scrapless nut, cold forging, cold heading, and cold extrusion qualities) for hot-rolled bars used in the production of solid or hollow shapes by means of severe cold plastic deformation, such as (but not limited to) upsetting, heading, forging, and forward or backward extrusion involving movement of metal by expansion and/or compression. Such processing normally involves special inspection standards and requires sound steel of special surface quality and uniform chemical composition. If steel of the type or chemical composition specified does not have adequate cold-forming characteristics in the as-rolled condition, a suitable heat treatment, such as annealing or spheroidize annealing, may be necessary. Axle Shaft Quality. Bars of axle shaft quality are produced for the manufacture of power-driven axle shafts for cars, trucks, and other vehicles. Because of their design or method of manufacture, these axles either are not machined all over or undergo less than the recommended amount of stock removal for proper cleanup of normal surface imperfections. Therefore, it is necessary to minimize the presence of injurious surface imperfections in bars of axle shaft quality through the use of special rolling practices, special billet and bar conditioning, and selective inspection. Cold-Shearing Quality. There are limits to the sizes of hot-rolled steel bars that can normally be cold sheared without specially controlled production procedures. When the cold shearing of larger bars is desirable, it is recommended that cold-shearing quality bars be ordered. Bars of this quality have characteristics that prevent cracking even in these larger sizes. Cold-shearing quality bars are not produced to specific requirements such as hardness, microstructure, shear life, or productivity. Maximum size (cross-sectional area) limitations for the cold shearing of hot-rolled steel bars without the specially controlled production procedures, and of cold-shearing quality bars, are given in the AISI manual that covers hot-rolled bars. If even larger bars are to be cold sheared, cold-shearing behavior can be further improved by suitable prior heat treatment. Structural quality is the descriptor for hot-rolled bars used in the construction of bridges and buildings by riveting, bolting, or welding and for general structural purposes. The general requirements for bars of this quality are given in ASTM A 6; individual ASTM specifications are listed in Table 3 . Additional qualities of carbon steel bars are available for specific requirements. Such qualities are related to application and processing. They include: • • • • • • • •

File quality Gun barrel quality Gun receiver quality Shell steel quality A Shell steel quality B Shell steel quality C Shell steel quality D Standard tube round quality

Alloy Steel Bars Hot-rolled alloy steel bars are commonly produced in the same size as special quality steel bars. Common size ranges have not been established for other shapes of hot-rolled alloy steel bar, including bar-size shapes, ovals, half-ovals, half-rounds, octagons, and special bar-size shapes. Hot-rolled alloy steel bars are covered by several ASTM specifications (Tables 3 and 4 ). Many of the alloys covered in these specifications are standard AISI-SAE grades (Table 5 ). Table 5 AISI-SAE grades of hot-rolled alloy steel bars in ASTM specifications ASTM specification

AISI-SAE grades

A 295

52100, 51100, 50100

A 304

All H grades except 4626H and 86B30H

A 322

All standard grades except 4032, 4042, 4135, 4422, 4427, 4617, 50B40, 5046, 5060, 5115, 5117, 50100, 8115, 86B45, 8650, 8660, 9310, and 94B15

A 434

By agreement

A 534

4023, 4118, 4320, 4620, 4720, 5120, 8620, E-3310, E-9310

A 535

3310, 4320, 4620, 4720, 4820, 52100, 52100 Mod. 1, 52100 Mod. 2, 52100 Mod. 3, 52100 Mod. 4, 8620, 9310

Hot-rolled alloy steel bars are also covered by several quality descriptors, which are discussed below. As with all quality descriptors, these descriptors differentiate bars on the basis of characteristic properties required to meet the particular conditions encountered during fabrication or use. Regular quality is the basic or standard quality for hot-rolled alloy steel bars, such as those covered by ASTM A 322. Steel for this quality are killed, are usually produced to fine grain size, and are melted to chemical composition limits. Bars of this


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quality are inspected, conditioned, and tested to meet the normal requirements for regular construction applications for which alloy steel is used. Axle Shaft Quality. Alloy steel bars of axle shaft quality are similar to carbon steel bars of the same quality (see the discussion of axle shaft quality bars in the section "Carbon Steel Bars for Specific Applications" in this article). Ball and roller bearing quality and bearing quality apply to alloy steel bars intended for antifriction bearings. These bars are usually made from steels of the AISI-SAE standard alloy carburizing grades and the AISI-SAE high-carbon chromium series. These steels can be produced in accordance with ASTM A 534, A 535, A 295, or A 485 (Table 4 ). Bearing quality steel bars require restricted melting and special teeming, heating, rolling, cooling, and conditioning practices to meet rigid quality standards. Steelmaking practices may include vacuum treatment. The foregoing requirements include thorough examination for internal imperfections by one or more of the following methods: macroetch testing, microscopic examination for nonmetallic inclusions, ultrasonic inspection, and fracture testing. It is not practical to furnish bearing quality steel bars in sizes exceeding 64,500 mm2 (100 in.2) in cross-sectional area to the same rigid requirements as those for bars in smaller sizes because of insufficient hot working in the larger bars. Usually, bars over 102 mm (4 in.) in thickness are forged to 102 mm (4 in.) square or smaller for testing. Cold-Shearing Quality. Alloy steel bars of cold-shearing quality are similar to carbon steel bars of the same quality (see the discussion of cold-shearing quality bars in the section "Carbon Steel Bars for Specific Applications" in this article). Cold-working quality, which replaces the older terminologies cold-heading quality and special cold-heading quality, is the descriptor for hot-rolled bars used in the production of solid or hollow shapes by means of severe cold plastic deformation, such as (but not limited to) upsetting, heading, forging, and forward or backward extrusion involving movement of metal by expansion and/or compression. Such processing normally involves special inspection standards and requires sound steel of special surface quality and uniform chemical composition. If steel of the type or chemical composition specified does not have adequate cold-forming characteristics in the as-rolled condition, a suitable heat treatment, such as annealing or spheroidize annealing, may be necessary. Aircraft quality and magnaflux quality are the descriptors used for alloy steel bars for critical or highly stressed parts of aircraft and for other similar or corresponding purposes involving additional stringent requirements such as magnetic particle inspection, additional discard, macroetch tests, and hardenability control. To meet these requirements, exacting steelmaking, rolling, and testing practices must be employed. These practices are designed to minimize detrimental inclusions and porosity. Phosphorus and sulfur are usually limited to 0.025% maximum each. Many parts for aircraft, missiles, and rockets require aircraft quality alloy steel bars. Magnetic particle testing as in AMS 2301 is sometimes specified for such applications. Some very critical aircraft, missile, and rocket applications require alloy steel bars of a quality attained only by vacuum melting or by an equivalent process. The requirements of AMS 2300 are sometimes specified for such applications. Aircraft quality alloy steel bars are ordinarily made to Aerospace Material Specifications published by the Society of Automotive Engineers. Typical examples of parts for aircraft engines and airframes made from bars covered by AMS specifications are given in Table 6 . Table 6 Specifications and grades of alloy steel bars for aircraft parts

Part Aileron, rudder, and, and elevator hinge pins Airframe parts (tubing, fittings, and braces)

AMS specification

AISI-SAE grade or approximate grade

6415

E4340

6370

4130

6280

8630

6382

4140

6322

8740

6415

E4340

Bearings

6440

E52100

Bolts, studs, and nuts

6322

8740

Connecting rods

6415

E4340

Crankcases

6342

9840

6382

4140

6322

8740

Crankshafts

6415

E4340

Gears and shafts

6415

E4340

6448

6150

6274

8620

6322

8740

Landing gears


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01 Sep 2005

6382

4140

6415

E4340

Propellers, spiders, hubs, and barrels

6415

E4340

Springs

6450

6150

Structural quality is the descriptor for hot-rolled bars used in the construction of bridges and buildings by riveting, bolting, or welding and for general structural purposes. The general requirements for bars of this quality are given in ASTM A 6; the only individual ASTM specification referred to in A 6 that pertains to alloy steel bars is A 710. Additional Qualities. The quality designations shown below apply to alloy steel bars intended for rifles, guns, shell, shot, and similar applications. They may involve requirements for amount of discard, macroetch testing, surface quality, or magnetic particle testing, as indicated in the product specification: • • • • • •

AP shot quality AP shot magnaflux quality Gun quality Rifle barrel quality Shell quality Shell magnaflux quality

High-Strength Low-Alloy Steel Bars In addition to the carbon steel and alloy steel bars of structural quality discussed in preceding sections of this article, ASTM A 6 also lists several specifications covering high-strength low-alloy (HSLA) steel bars of structural quality (Table 3 ). High-strength low-alloy steel bars are also covered in SAE J 1442. Bars of these steels offer higher strength than that of carbon steel bars and are frequently selected for applications in which weight saving is important. They also offer increased durability, and many offer increased resistance to atmospheric corrosion. Additional information on HSLA steels is available in the articles "High-Strength Structural and High-Strength Low-Alloy Steels," "High-Strength Low-Alloy Steel Forgings" and "Bulk Formability of Steels" in this Volume. Microalloyed steel bars constitute a class of special quality carbon steels to which small amounts of alloying elements such as vanadium, niobium, or titanium have been added. Microalloyed steels in the as-hot-rolled condition are capable of developing strengths higher than those of the base carbon grades through precipitation hardening. In some cases, strength properties comparable to those of the quenched and tempered base grade can be attained. These steels are finding increased application in shafting and automotive forgings.

Concrete-Reinforcing Bars Concrete-reinforcing bars are available as either plain rounds or deformed rounds. Deformed reinforcing bars are used almost exclusively in the construction industry to furnish tensile strength to concrete structures. The surface of the deformed bar is provided with lugs, or protrusions, which inhibit longitudinal movement relative to the surrounding concrete. The lugs are hot formed in the final roll pass by passing the bars between rolls into which patterns have been cut. Plain reinforcing bars are used more often for dowels, spirals, structural ties, and supports than as substitutes for deformed bars. Concrete-reinforcing bars are supplied either straight and cut to proper length, or bent or curved in accordance with plans and specifications. Grades. Deformed and plain concrete-reinforcing bars rolled from billet steel are produced to the requirements of ASTM A 615 or A 706. For special applications that require deformed bars with a combination of strength, weldability, ductility, and improved bending properties, ASTM A 706 is specified, which is an HSLA steel. Deformed and plain concrete-reinforcing bars are also available rolled from railroad rails (ASTM A 616) and from axles for railroad cars (ASTM A 617), Specification ASTM A 722 covers deformed and plain uncoated high-strength steel bars for prestressing concrete structures. Sizes. Numbers indicating sizes of reinforcing bars correspond to nominal bar diameter in eighths of an inch for sizes 3 through 8; this relationship is approximate for sizes 9, 10, 11, 14, and 18. The nominal values for bar diameter, cross-sectional area, and weight per unit length corresponding to these size numbers are given in Table 7 . The nominal cross-sectional area and the nominal diameter of a deformed bar are the same as those of a plain bar of equal weight per foot. Table 7 Dimensions of deformed and plain concrete-reinforcing bars of standard sizes Nominal diameter

Bar size

mm

3

9.52

4

12.70

Crosssectional area

Nominal weight

in.

mm2

in.2

kg/m

lb/ft

0.375

71

0.11

0.560

0.376

0.500

129

0.20

0.994

0.668


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01 Sep 2005

5

15.88

0.625

200

0.31

1.552

1.043

6

19.05

0.750

284

0.44

2.235

1.502

7

22.22

0.875

387

0.60

3.042

2.044

8

25.40

1.000

510

0.79

3.973

2.670

9

28.65

1.128

645

1.00

5.059

3.400

10

32.26

1.270

819

1.27

6.403

4.303

11

35.81

1.410

1006

1.56

7.906

5.313

14

43.00

1.693

1452

2.25

11.384

7.65

18

57.33

2.257

2581

4.00

20.238

13.60

Structural Shapes Structural shapes, as stated previously, are flanged shapes 76 mm (3 in.) and greater in at least one cross-sectional dimension (smaller shapes are referred to as bar-size shapes) and are used in the construction of structures such as bridges, buildings, ships, and railroad cars. Included in this product category are regular structural shapes (see ASTM A 6), such as standard beams, wide-flange beams, columns, light beams, joists, stanchions and bearing piles, and certain tees, along with special structural shapes, which are those designed for specialized applications and that have dimensions and/or values of weight per foot that do not conform to regular shapes. Bar-size structural shapes (angles, channels, tees, and zees with greatest cross-sectional dimension less than 76 mm, or 3 in.) are considered to be in the merchant quality bar category rather than the structural shape category. The common method of designating sizes of structural shapes is as follows: • Beams and channels: By depth of cross section and weight per foot. • Angles: By length of legs and thickness in fractions of an inch or, more commonly, by length of legs and weight per foot. The longer leg of an unequal angle is commonly stated first • Tees: By width of flange, overall depth of stem, and weight per foot, in that order • Zees: By depth, width of flanges, and thickness in fractions of an inch or by depth, flange width, and weight per foot • Wide-flange shapes: By depth, width across flange, and weight per foot, in that order Most structural shapes are produced to meet specific standard specifications, such as those listed in Table 3 . Structural shapes are generally furnished to chemical composition limits and mechanical property requirements. Special requirements are sometimes specified for structural shapes to adapt them to conditions they will encounter during fabrication or service. These requirements may include specific deoxidation practices, additional mechanical tests, or nondestructive testing.

Special Shapes Special shapes are hot-rolled steel shapes made with cross-sectional configurations uniquely suited to specific applications. Examples of custom-designed shapes are track shoes for tractors or tanks and sign-post standards. The only type of standard shape in high production that falls in this classification is rail. Railroad rails of the standard American tee rail shape are produced from carbon steel to the dimensional, chemical, and other requirements of the American Railway Engineering Association (AREA). The sizes of railroad rails are designated in pounds per yard of length; rails are furnished in 40 to 64 kg (90 to 140 lb) sizes. The most common sizes are 52, 60, 62, and 64 kg (115, 132, 136, and 140 lb). The ordinary length of railroad rails is 12 m (39ft). Carbon steel tee rails for railway track are covered by ASTM A 1; rail-joint bars and tie plates are covered in ASTM A 3, A 4, A 5, A 49, A 67, and A 241. Light rails are available for light duty, such as in mines and amusement park rides, in sizes from 6.8 to 39 kg (15 to 85 lb). Light rails are covered by specifications of the American Society of Civil Engineers (ASCE). Crane rails generally have heavier heads and webs than those of railroad rails in order to withstand the heavy loads imposed on them in service. Crane rails in sizes from 18 to 79 kg (40 to 175 lb) are furnished to ASCE, ASTM, and producers' specifications. REFERENCES 1. Alloy, Carbon and High Strength Low Alloy Steels: Semifinished for Forging; Hot Rolled Bars, Cold Finished Bars; Hot Rolled Deformed and Plain Concrete Reinforcing Bars, AISI Steel Products Manual, American Iron and Steel Institute, 1986 2. Materials, Vol 1, 1989 SAE Handbook, Society of Automotive Engineers, 1989


Page 412


Cold-Finished Steel Bars

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Carbon and Low-Alloy Steels Cold-Finished Steel Bars Revised by the ASM Committee on Cold-Finished Bars* [*K. M. Shupe, Bliss & Laughlin Steel Company; Richard B. Smith, Stanadyne Western Steel; Steve Slavonic, Teledyne Columbia-Summerill; B. F. Leighton, Canadian Drawn Steel Company; W. Gismondi, Union Drawn Steel Company, Ltd.; John R Stubbles, LTV Steel Company; Kurt W. Boehm, Nucor Steel; Donald M. Keane, LaSalle Steel Company] COLD-FINISHED STEEL BARS are carbon and alloy steel bar products (round, square, hexagonal, flat, or special shapes) that are produced by cold finishing previous hot-wrought bars by means of cold drawing, cold forming, turning, grinding, or polishing (singly or in combination) to yield straight lengths or coils that are uniform throughout their length. Not covered in this article are flat-rolled products such as sheet, strip, or plate, which are normally cold finished by cold rolling, or cold-drawn tubular products. Cold-finished bars fall into five classifications: • • • • •

Cold-drawn bars Turned and polished (after cold drawn or hot roll) bars Cold-drawn, ground, and polished (after cold draw) bars Turned, ground, and polished bars Cold-drawn, turned, ground, and polished bars

Cold-drawn bars represent the largest tonnage production and are widely used in the mass production of machined and other parts. They have attractive combinations of mechanical and dimensional properties. Turned and polished bars have the mechanical properties of hot-rolled products but have greatly improved surface finish and dimensional accuracy. These bars are available in sizes lager than those that can be cold drawn. Turned bars are defect and decarb free. Cold-drawn, ground, and polished bars have the increased machinability, tensile strength, and yield strength of cold-drawn bars together with very close size tolerances. However, cold-drawn, ground, and polished bars are not guaranteed to be defect free. Turned, ground, and polished bars have superior surface finish, dimensional accuracy, and straightness. These bars find application in precision shafting and in plating, where such factors are of primary importance. Cold-drawn, turned, ground, and polished bars have improved mechanical properties, close size tolerances, and a surface free of imperfections.

Bar Sizes Cold-finished steel bars are available in a wide variety of sizes and cross-sectional shapes. Normally, they are furnished in straight lengths, but in some sizes and cross sections they may be furnished in coils. Cold-finished steel bars are available with nominal dimensions designated in either inches or millimeters. Cold-finished product is available in standard size increments, which vary by size range. Special sizes can be negotiated depending on hot mill increments and cold-finish tooling. The sizes in which they are commonly available in bar and coil form are given in Table 1 . Table 1 Common commercially available sizes of cold-finished steel bars and coils Bars(a) Minimum thickness or diameter

Maximum thickness or diameter

Configuration

mm

in.

mm

in.

Round

3.2

0.125

305

12


Size increments mm

in.

Normal length m

0.8−2 32nds to 1 in., 16ths 3.0−3.7 to 3 in., 8ths to 6 or 5 in. 6.1−7.3 1.6−7

Coils(b), sizes

ft

mm

in.

10−12 or 20−24

≤25

≤1

Page 413



01 Sep 2005

5 3.2−1 52 Square

3.2

0.125

152

6

1.6−3 8 3.2−7 0

16ths to 11=2 in., 8ths to 23=4 in.

3.0−3.7

10−12

≤16

≤5=8

Hexagonal

3.2

0.125

102

4

1.6−5 0 6.4−1 02

16ths to 2 in., 4ths to 4 in.

3.0−3.7

10−12

≤16

≤5=8

3 thick × 1.6−1 16ths to 11=16 in., 3.0−3.7 8ths 7 145=8 3 wide 3.2−4 to 1 =4 in., 4ths to 3 in 4 6.4−7 6

10−12

≤14.3 × 15.9(c)

≤9=16 × 5 =8(c)

Flat

3.2 thick 0.125 thick × 6.4 × 0.25 wide wide

76 × 371

(a) Ref 1. (b) Ref 2. (c) Or other sections having cross-sectional areas ≤194 mm2 (≤0.30 in.2)

Product Types In the manufacture of cold-finished bars, the steel is first hot rolled oversize to appropriate shape and is then subjected to mechanical operations (other than those intended primarily for scale removal) that affect is machinability, straightness, and end-cut properties. The two common methods of cold finishing bars are: • Removal of surface material by turning or grinding, singly or in combination • Drawing the material through a die of suitable configuration Pickling or blasting to remove scale may precede turning or grinding and must always precede drawing. For bar products, cold rolling has been almost superseded by cold drawing. Nevertheless, cold-finished bars and special shapes are sometimes incorrectly described as cold rolled. Commercial Grades. Any grade of carbon or alloy steel that can be hot rolled can also be cold finished. The choice of grade is based on the attainable cold-finished and/or hardenability and tempering characteristics necessary to obtain the required mechanical properties. Production methods vary widely among cold-finished cold-drawn suppliers. For example, one supplier currently anneals and cold draws grades 1070, 1090, and 5160, and in the future plans to do the same with grade 9254. Grade 1070 is a high-volume item, and cold drawing is required for precision sizing and subsequent nondestructive testing of the bar, using a rotating-probe eddy current device (see the articles "Eddy Current Inspection," "Remote-Field Eddy Current Inspection," and "Steel Bar, Wire, and Billets" in Nondestructive Evaluation and Quality Control, Volume 17 of ASM Handbook, formerly 9th Edition Metals Handbook) for detecting surface seams. Cold drawing is also necessary because the smallest hot-rolled size typically available for some applications is not small enough for customer use. Thus, a supplier whose smallest hot-rolled bar size is 11.1 mm (0.437 in.) cold draws this diameter to as small as 9.98 mm (0.393 in.). Carbon steels containing more than 0.55% C must be annealed prior to being cold drawn so that the hardness will be sufficiently low to facilitate the cold-drawing operation. For carbon steels containing up to 0.65% C, this will normally be a lamellar pearlitic anneal; for carbon steels containing more than 0.65% C, a spheroidize anneal is required. The type of structure required is normally reached by agreement between the steel producer and the customer. Alloy steels containing more than 0.38% C are usually annealed before cold drawing. Machined Bars. Bar products that are cold finished by stock removal can be: • • • • •

Turned and polished Turned, ground, and polished Cold drawn, ground, and polished Cold drawn, turned, and polished Cold drawn, turned, ground, and polished

Turning is done in special machines with cutting tools mounted in rotating heads, thus eliminating the problem of having to support long bars as in a lathe. Grinding is done in centerless machines. Polishing can be done in a roll straightener of the crossed-axis (Medart) type with polished rolls to provide a smooth finish. Polishing by grinding with an organic wheel or with a belt is of increasing interest (see the article "Grinding Equipment and Processes" in Machining, Volume 16 of ASM Handbook, formerly 9th Edition Metals Handbook) because it is cost effective to grind and polish the bars on the same machine simply by


Page 414



01 Sep 2005

using grinding wheels or belts of different grit size. Grinding produces a smoother finish than turning; polishing improves the surface produced by either technique. Turned, ground, and polished rounds represent the highest degree of overall accuracy, concentricity, straightness, and surface perfection attainable in commercial practice (Ref 3). The surface finish desired is specified by using the process names given above because the industry has not developed standard numerical values for roughness, such as microinch or root mean square (rms) numbers. However, surface finish with respect to rms (root mean square deviation from the mean surface) as determined with a profilometer can be negotiated between the producer and a customer. This could be done for such critical-finish applications as turned and polished bars used to produce shafting as well as stock used to produce machined parts of which a superior finish is required on surfaces not machined. The published range of diameters both for turned and for turned and ground bars is 13 to 229 mm (1=2to 9 in.) inclusive; for cold-drawn and ground bars, it is 3.2 to 102 mm (1=8to 4 in.) inclusive. These are composites of size ranges throughout the industry; an individual producer may be unable to furnish a full range of sizes. For example, one well-known producer supplies turned rounds from 13 to 229 mm (1=2to 9 in.), another from 29 to 203 mm (11=8to 8 in.)⎯all finished sizes. Yet another producer supplies sizes up to and including 152 mm (6 in.) that are turned on special turning machines and ground on centerless grinders; larger sizes are lathe turned and ground on centers. Because turning and grinding do not alter the mechanical properties of the hot-rolled bar, this product can be machined asymmetrically with practically no danger of warpage (Ref 3). Stock removal is usually dependent on American Iron and Steel Institute (AISI) seam allowances (Ref 2). Stock removal in turning, or turning and grinding, measured on the diameter, is normally 1.6 mm (1=16in.) for sizes up to 38 mm (11=2in.), 3.2 mm (1=8in.) for the 38 to 76 mm (11=2to 3 in.) range, 4.8 mm (3=16in.) for the 76 to 127 mm (3 to 5 in.) range, and 6.4 mm (1=4in.) for 127 mm (5 in.) diameter and larger. Cold-drawn round bars are available in a range of diameters from 3.2 to 152 mm (1=8to 6 in.). The maximum diameters available from individual producers, however, may vary from 76 to 152 mm (3 to 6 in.). The reduction in diameter in cold drawing, called draft, is commonly 0.79 mm (1=32in.) for finished sizes up to 9.5 mm (3=8in.) and 1.6 mm (1=16in.) for sizes over 9.5 mm (3=8in.). Some special processes use heavier drafts followed by stress relieving. One producer employs heavy drafting at elevated temperature. With this exception, drawing operations are begun with the material at room temperature to start, and the only elevated temperature involved is that developed in the bar as a result of drawing; this temperature rise is small and of little significance. Originally, cold finishing, whether by turning or by cold rolling, was employed only for sizing to produce a bar with closer dimensional tolerances and a smoother surface. As cold-finished bar products were developed and improved, increased attention was paid to the substantial enhancement of mechanical properties that could be obtained by cold working. This additional advantage is now more fully appreciated, as evidenced by the fact that increased mechanical properties are an important consideration in about 40% of the applications. In approximately half of these applications, or 20% of the total, cold drawing is used only to increase strength; in the other 20%, close tolerances and better surface finish are desired in addition to increased strength. As-rolled microalloyed high-strength low-alloy (HSLA) steels or microalloyed HSLA steels in various combinations of controlled drafting and furnace treatment provide an extension of property attainment. A high percentage of free-machining steels are cold drawn for the combination of size accuracy and improved machinability. Recent developments in microalloyed steels provide hot-rolled turned bars, under certain circumstances, having mechanical properties similar to cold-drawn nonmicroalloyed steels. An appreciable fraction of all applications of cold finishing to carbon steel bars utilizes cold drawing to improve mechanical properties. For alloy steel, however, cold finishing is commonly used to improve surface finish and dimensional accuracy, and not for additional mechanical strength. When additional mechanical strength is desired, alloy steel bars may be heat treated (quenched and tempered) and then cold drawn and stress relieved. Elevated-temperature or warm-drawn steels are also available with increased mechanical strength and improved machinability. Heavily drafted and strain-tempered carbon and alloy steels subjected to induction hardening of the surface provide many additional property combinations. The extra cost of using alloy steel in cold-finished bars can be justified only when heat treatment (quenching and tempering) is necessary for meeting the required strength level. Because work-hardening effects are removed during heating prior to quenching, the benefit of increased mechanical strength due to cold finishing is eliminated from the finished product. Turning Versus Cold Drawing. Basic differences exist between bars finished by turning and those finished by cold drawing. First, it is obvious that turning and centerless grinding are applicable only to round bars, while drawing can be applied to a variety of shapes. Drawing, therefore, is more versatile than turning. Second, there is a difference in the number and severity of the surface imperfections that may be present. Because stock is removed in turning and grinding, shallow surface imperfections and decarburization may be completely eliminated. When material is drawn, stock is only displaced, and surface imperfections are only reduced in depth (in the ratio of the change in bar diameter or section thickness). The length of these imperfections may be slightly increased because in the drawing operation an increase in length accompanies the reduction in cross section. Cold-drawn bars can approach the freedom from surface imperfections obtained in turned or turned and ground bars if the hot-rolled bars from which they are produced are rolled from specially conditioned billets. Quality conditions such as cold-working quality are available from producers of hot-rolled bars. The depth limits of the surface imperfections are as agreed to between the producer and the customer. However, if maximum freedom from surface imperfections is the controlling factor, turned bars have an advantage.


Page 415



01 Sep 2005

Different size tolerances are applicable to cold-finished products, depending on shape, carbon content, and heat treatment. Listed in Tables 2 , 3 , and 4 are the tolerances for cold-finished carbon and alloy steel bars published in ASTM A 29. These tables include cold-drawn bars; turned and polished rounds; cold-drawn, ground, and polished rounds; and turned, ground, and polished rounds. From the data in Tables 2 , 3 , and 4 , certain generalizations can be stated. The tolerances for cold-drawn and for turned and polished rounds, for example, are the same for sizes up to and including 102 mm (4 in.). There are differences, however, between the tolerances that apply to carbon steel and those that apply to alloy steels. Tolerances for several finishes also vary with certain levels of carbon content. Broader tolerances are applicable to bars that have been heat treated before cold finishing. In contrast, tolerances are closer when bars are ground, and these tolerances are independent of carbon content. Table 2 Size tolerances for cold-finished carbon steel bars, cold drawn or turned and polished This table includes tolerances for bars that have been annealed, spheroidize annealed, normalized, normalized and tempered, or quenched and tempered before cold finishing. This table does not include tolerances for bars that are annealed, spheroidize annealed, normalized, normalized and tempered, or quenched and tempered after cold finishing; the producer should be consulted for tolerances for such bars. Size tolerance Maximum carbon (C) range, %

C ≤ 0.28

Size mm

in.

mm

in.

0.28 < C ≤ 0.55 mm

in.

C ≤ 0.55 including stress relief or annealed after cold finishing

C > 0.55

All grades quenched and tempered or normalized before cold finishing

mm

in.

mm

in.

mm

in.

Rounds⎯cold drawn (to 102 mm, or 4 in., in size) or turned and polished −0.05 −0.002

−0.08

−0.003

−0.10

−0.004

−0.13

−0.005

−0.13

−0.005

>38−64 inclusive

1

>1 =2−2 =2 inclusive

−0.08 −0.003

−0.10

−0.004

−0.13

−0.005

−0.15

−0.006

−0.15

−0.006

>64−102 inclusive

>21=2−4 inclusive

−0.10 −0.004

−0.13

−0.005

−0.15

−0.006

−0.18

−0.007

−0.18

−0.007

>102−152 inclusive

>4−6 inclusive

−0.13 −0.005

−0.15

−0.006

−0.18

−0.007

−0.20

−0.008

−0.20

−0.008


>6−8 inclusive

−0.15 −0.006

−0.18

−0.007

−0.20

−0.008

−0.23

−0.009

−0.23

−0.009


>8−9 inclusive

−0.18 −0.007

−0.20

−0.008

−0.23

−0.009

−0.25

−0.010

−0.25

−0.010

To 38 inclusive To 1 1=2inclusive 1

Hexagons⎯cold drawn −0.05 −0.002

−0.08

−0.003

−0.10

−0.004

−0.15

−0.006

−0.15

−0.006

>19−38 inclusive

>3=4−11=2 inclusive

−0.08 −0.003

−0.10

−0.004

−0.13

−0.005

−0.18

−0.007

−0.18

−0.007

>38−64 inclusive


−0.10 −0.004

−0.13

−0.005

−0.15

−0.006

−0.20

−0.008

−0.20

−0.008

64−80 inclusive >21=2−31=8 inclusive

−0.13 −0.005

−0.15

−0.006

−0.18

−0.007

−0.23

−0.009

−0.23

−0.009

−0.13 −0.005

−0.15

−0.006

...

...

...

...

...

...

−0.05 −0.002

−0.10

−0.004

−0.13

−0.005

−0.18

−0.007

−0.18

−0.007

To 19 inclusive To 3=4inclusive

>80−102 inclusive

>3 =8−4 inclusive 1

Squares⎯cold drawn(a) To 19 inclusive To 3=4 inclusive >19−38 inclusive

> =4−1 =2 inclusive

−0.08 −0.003

−0.13

−0.005

−0.15

−0.006

−0.20

−0.008

−0.20

−0.008

>38−64 inclusive


−0.10 −0.004

−0.15

−0.006

−0.18

−0.007

−0.23

−0.009

−0.23

−0.009

>64−102 inclusive

>21=2−4 inclusive

−0.15 −0.006

−0.20

−0.008

−0.23

−0.009

−0.28

−0.011

−0.28

−0.011


>4−5 inclusive

−0.25 −0.010

...

...

...

...

...

...

...

...


>5−6 inclusive

−0.36 −0.014

...

...

...

...

...

...

...

...

3

1

Flats⎯cold drawn(a)(b)


Page 416



01 Sep 2005

−0.08 −0.003

−0.10

−0.004

−0.15

−0.006

−0.20

−0.008

−0.20

−0.008

>19−38 inclusive


−0.10 −0.004

−0.13

−0.005

−0.20

−0.008

−0.25

−0.010

−0.25

−0.010

>38−75 inclusive

>11=2−3 inclusive

−0.13 −0.005

−0.15

−0.006

−0.25

−0.010

−0.30

−0.012

−0.30

−0.012

>75−102 inclusive

>3−4 inclusive

−0.15 −0.006

−0.20

−0.008

−0.28

−0.011

−0.40

−0.016

−0.40

−0.016


>4−6 inclusive

−0.20 −0.008

−0.25

−0.010

−0.30

−0.012

−0.50

−0.020

−0.50

−0.020

To 19 inclusive To 3=4 inclusive

>152 >6 ... ... ... ... ... ... −0.33 −0.013 −0.38 −0.015 (a) Tolerances can be ordered all plus, or distributed plus and minus with the sum equivalent to the tolerances listed. (b) Width governs the tolerance for both width and thickness of flats, for example, when the maximum of carbon range is 0.28% or less for a flat 50 mm (2 in.) wide and 25 mm (1 in.) thick. The width tolerance is 0.13 mm (0.005 in.), and the thickness is the same, nearly 0.13 mm (0.005 in.). Source: Ref 4

Table 3 Size tolerances for cold-finished alloy steel bars, cold drawn or turned and polished This table includes tolerances for bars that have been annealed, spheroidize annealed, normalized, normalized and tempered, or quenched and tempered before cold finishing. This table does not include tolerances for bars that are annealed, spheroidize annealed, normalized, normalized and tempered, or quenched and tempered after cold finishing; the producer should be consulted for tolerances for such bars. Size tolerance Maximum carbon (C) range, %

C ≤ 0.28

Size mm

in.

mm

in.

0.28 < C ≤ 0.55 mm

in.

All carbons C ≤ 0.55 quenched including stress C > 0.55 with or and tempered without stress (heat treated) or relief or relieving or normalized and annealed annealing after tempered before after cold cold finishing cold finishing finishing mm

in.

mm

in.

mm

in.

Rounds⎯cold drawn (to 102 mm, or 4 in., in size) or turned and polished In coils: To 25 inclusive

To 1 inclusive

0.05

0.002

0.08

0.003

0.10

0.004

0.13

0.005

0.13

0.005

Cut lengths: To 38 inclusive

>To 11=2 inclusive

0.08

0.003

0.10

0.004

0.13

0.005

0.15

0.006

0.15

0.006

>38−64 inclusive


0.10

0.004

0.13

0.005

0.15

0.006

0.18

0.007

0.18

0.007

>64−102 inclusive

>2 =2−4 inclusive

0.13

0.005

0.15

0.006

0.18

0.007

0.20

0.008

0.20

0.008


>4−6 inclusive

0.15

0.006

0.18

0.007

0.20

0.008

0.23

0.009

0.23

0.009


>6−8 inclusive

0.18

0.007

0.20

0.008

0.23

0.009

0.25

0.010

0.25

0.010


>8−9 inclusive

0.20

0.008

0.23

0.009

0.25

0.010

0.28

0.011

0.28

0.011

1

Hexagons⎯cold drawn To 19 inclusive

To 3=4inclusive

0.08

0.003

0.10

0.004

0.13

0.005

0.18

0.007

0.18

0.007

>19−38 inclusive


0.10

0.004

0.13

0.005

0.15

0.006

0.20

0.008

0.20

0.008

>38−64 inclusive


0.13

0.005

0.15

0.006

0.18

0.007

0.23

0.009

0.23

0.009

64−79 inclusive

1


0.15

0.006

0.18

0.007

0.20

0.008

0.25

0.010

0.25

0.010

>79−102 inclusive

>3 =8−4 inclusive

0.15

0.006

...

...

...

...

...

...

...

...

To 19 inclusive

To 3=4 inclusive

0.08

0.003

0.13

0.005

0.15

0.006

0.20

0.008

0.20

0.008

>19−38 inclusive


0.10

0.004

0.15

0.006

0.18

0.007

0.23

0.009

0.23

0.009

3

1

1 1

Squares⎯cold drawn 3

1

>38−64 inclusive

1


0.13

0.005

0.18

0.007

0.20

0.008

0.25

0.010

0.25

0.010

>64−102 inclusive

>2 =2−4 inclusive

0.18

0.007

0.23

0.009

0.25

0.010

0.30

0.012

0.30

0.012


>4−5 inclusive

0.28

0.011

0.23

...

0.25

...

0.30

...

0.30

...

1 1

Flats⎯cold drawn(a) To 19 inclusive

To 3=4 inclusive

0.10

0.004

0.13

0.005

0.18

0.007

0.23

0.009

0.23

0.009

>19−38 inclusive


0.13

0.005

0.15

0.006

0.23

0.009

0.28

0.011

0.28

0.011


Page 417



01 Sep 2005

>38−76 inclusive

>11=2−3 inclusive

0.15

0.006

0.18

0.007

0.28

0.011

0.33

0.013

0.33

0.013

>76−102 inclusive

>3−4 inclusive

0.18

0.007

0.23

0.009

0.30

0.012

0.43

0.017

0.43

0.017


>4−6 inclusive

0.23

0.009

0.28

0.011

0.33

0.013

0.52

0.021

0.52

0.021

>152 >6 0.36 0.014 ... ... ... ... ... ... ... ... (a) Width governs the tolerance for both width and thickness of flats, for example, when the maximum of carbon range is 0.28% or less for a flat 50 mm (2 in.) wide and 25 mm (1 in.) thick. The width tolerance is 0.13 mm (0.005 in.), and the thickness is the same, nearly 0.13 mm (0.005 in.). Source: Ref 4

Table 4 Size tolerances for cold-finished carbon and alloy steel round bars cold drawn, ground, and polished or turned, ground, and polished Size Cold drawn, ground, and polished mm

Tolerances from specified size

Turned, ground, and polished

mm

in.

To 38 incl

To 11=2 incl

in. To 38 incl

mm To 11=2 incl

−0.03

−0.001

>38−64 incl

>11=2−21=2 excl

>38−64 incl

>11=2−21=2 excl

−0.04

−0.0015

>64−76 incl

≥2 =2−3 incl

>64−76 incl

≥2 =2−3 incl

−0.05

−0.002

>76−102 incl

>3−4 incl

>76−102 incl

>3−4 incl

−0.08

−0.003

>102−152 incl

>4−6 incl

−0.10(a)

−0.004(a)

1

...

...

in.

1

... ... >152 >6 −0.13(a) −0.005(a) incl, inclusive; excl, exclusive.(a) For nonresulfurized steels (steels specified to maximum sulfur limits under 0.08% or for steels thermally treated, the tolerance is increased by 0.03 mm (0.001 in.). Source: Ref 4

In addition to the size-tolerance requirements for all cold-finished steel bars, straightness tolerances are also of major importance for bars intended for use in automatic screw machines. Table 5 (also from ASTM A 29) details the straightness requirements for rounds, squares, hexagons, and octagons, which are the same for both carbon and alloy steel bars. As indicated in Table 6 , special provisions are also made for bars subject to magnetic particle inspection. Table 5 Straightness tolerances for cold-finished carbon and alloy steel bars All grades quenched and tempered or normalized and tempered to ≤HB 302 before cold finishing; all grades stress relieved or annealed after cold finishing. Straightness tolerances are not applicable to bars having Brinell hardness exceeding 302. The tolerances are based on the following method of measuring straightness. Departure from straightness is measured by placing the bar on a level table so that the arc or departure from straightness is horizontal, and the depth of the arc is measured with a feeler gage and a straightedge. It should be recognized that straightness is a perishable quality and may be altered by mishandling. The preservation of straightness in cold-finished bars requires the utmost care in subsequent handling. Specific straightness tolerances are sometimes required for carbon and alloy steels, in which case the purchaser should inform the manufacturer of the straightness tolerances and the methods to be used in checking the straightness. Straightness tolerances (maximum deviation) from straightness in any 3 m (10 ft) portion of the bar Carbon range >0.28% and all grades thermally treated

Carbon range, ≤0.28%

Size mm

Length in.

m

Squares, hexagons, and octagons

Rounds ft

mm

in. 1 =8

36.5−49.2 incl

>1 =16 −1 =16 incl

2.29

0.090

>49.2−61.9 incl

>115=16−27=16 incl

3.18

0.125

>61.9−85.7 incl

>2 =16 −3 =8 incl

3.96

0.156

15

7

15

7

3

4.75 0.187 >85.7−111 incl >3 =8−4 =8 incl incl, inclusive.(a) For example, the minimum reduction in diameter of rounds is twice the minimum stock removal from the surface. Source: Ref 2 3

3

Product Quality Descriptors The term quality relates to the suitability of a mill product to become an acceptable part. When used to identify cold-finished steel bars, the various quality descriptors are indicative of many characteristics, such as degree of internal soundness, relative uniformity of chemical composition, and relative freedom from detrimental surface imperfections. Because of the characteristic surface finish of cold-drawn bars, close visual inspection cannot identify detrimental surface imperfections. Therefore, for applications that do not allow surface imperfections on the finished surfaces of standard quality cold-drawn carbon steel bars and regular quality cold-drawn alloy steel bars, the user should recognize that some stock removal is necessary to eliminate such imperfections as seams. The recommended stock removal per side for all nonresulfurized grades is 0.025 mm (0.001 in.) per 1.6 mm (1=16in.) of cross section, or 0.25 mm (0.010 in.), whichever is greater. For example, for a 25 mm (1 in.) bar, recommended stock removal is 0.41 mm (0.016 in.) per side. For the resulfurized grades, recommended stock removal is 0.038 mm (0.0015 in.) per 1.6 mm (1=16in.), or 0.38 mm (0.015 in.), whichever is greater. Therefore, for a 25 mm (1 in.) bar, recommended stock removal is 0.61 mm (0.024 in.) per side. Occasionally, some bars in a shipment may have imperfections that exceed the recommended stock removal limits. Therefore, for critical applications, inspection of finished parts is recommended, or more restrictive quality and/or additional inspection methods can be specified by agreement of both supplier and customer. To minimize pitting, the recommended stock removal per side for cold-drawn bars that are to be decorative chromium plated is as follows: Size, mm (in.)

Stock removal per side, mm (in.)

Through 7.9 (5=16 )

0.15 (0.006)

Over 7.9 (5=16 ) through 11.1 (7=16 )

0.20 (0.008)

Over 11.1 (7=16 )

0.25 (0.010)

Carbon Steel Quality Descriptors Standard quality is the descriptor applied to the basic quality level to which cold-finished carbon steel bars are produced. Standard quality cold-finished bars are produced from hot-rolled carbon steel of special quality (the standard quality for hot-rolled bars for cold finishing). Steel bars of standard quality must be free from visible pipe and excessive chemical segregation. They may contain surface imperfections. In general, the size of surface imperfections increases with bar size. Restrictive requirement quality A (RRA) incorporates all the features of standard quality carbon steel bars described above, plus any one of the following restrictive requirements. Special surface bars are produced with special surface preparation to minimize the frequency and size of seams and other surface imperfections. These bars are used for applications in which machining allowances do not allow sufficient surface removal to clean up the detrimental imperfections that occur in standard quality bars. Special internal soundness bars have greater freedom from chemical segregation and porosity than standard quality bars. Special hardenability bars are produced to hardenability requirements other than those of standard H-steels. Cold-finished carbon steel bars are also produced to inclusion ratings as determined by standard nonmetallic inclusion testing. Restrictive requirement quality B (RRB) incorporates all the features of standard quality carbon steel bars, plus any one of the following. Special discard is specified when minimized chemical segregation, special steel cleanliness, or internal soundness requirements dictate that the product be selected from certain positions in the ingot. Minimized decarburization is specified whenever decarburization is important, as in heat treating for surface hardness requirements.


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Single restrictions other than those noted above, such as special chemical limitations, special processing techniques, and other special characteristics not previously anticipated, are also covered by this quality level. Multiple restrictive requirement quality (MRR) applies when two or more of the above-described restrictive requirements are involved. Cold-forging quality A and cold-extrusion quality A apply to cold-finished carbon steel bars used in the production of solid or hollow shapes by means of cold plastic deformation involving the movement of metal by compression with no expansion of the surface and not requiring special inspection standards. For an individual application, if the type of steel or chemical composition specified does not provide adequate cold-forming characteristics in the as-drawn condition, a suitable heat treatment to provide proper hardness or microstructure may be necessary. Cold-heading quality, cold-extrusion quality B, cold-upsetting quality, and cold-expansion quality apply to cold-finished carbon steel bars used in production of solid or hollow shapes by means of severe cold plastic deformation by cold heading, cold extrusion, cold upsetting, or cold expansion involving movement of metal by expansion and/or compression. Such bars are obtained from steel produced by closely controlled steelmaking practices and are subject to special inspection standards for internal soundness and surface quality and uniform chemical composition. For grades of steel with a maximum specified carbon content of 0.30% or more, an anneal or spheroidize anneal heat treatment may be required to obtain the proper hardness and microstructure for cold working. Restrictive cold-working quality applies to cold-finished carbon steel bars used in the production of solid or hollow shapes by means of very severe cold plastic deformation involving cold working by expansion and/or compression. This degree of cold working normally involves restrictive inspection standards and requires steel that is exceptionally sound, of uniform chemical composition, and virtually free of detrimental surface imperfections. Such severe cold-forming operations normally require suitable heat treatment to obtain proper hardness and microstructure for cold working. Other Carbon Steel Qualities. The quality descriptors listed below are some of those that apply to cold-finished carbon steel bars intended for specific requirements and applications. They may have requirements for surface quality, amount of discard, macroetch tests, mechanical properties, or chemical uniformity as indicated in product specifications: • • • • •

Axle shaft quality Shell steel quality A Shell steel quality C Rifle barrel quality Spark plug quality

Alloy Steel Quality Descriptors Regular quality is the descriptor applied to the basic, or standard, quality level to which cold-finished alloy steel bars are produced. Steels for this quality are killed and are usually produced to a fine grain size. They are melted to chemical ranges and limits and are inspected and tested to meet normal requirements for regular constructional alloy steel applications. Regular quality cold-finished alloy steel bars may contain surface imperfections to the depths mentioned in the opening paragraphs of the section "Product Quality Descriptors" in this article. In general, the size of detrimental surface imperfections increases with bar size. Cold-heading quality applies to cold-finished alloy steel bars intended for applications involving cold plastic deformation by such operations as upsetting, heading, or forging. Bars are supplied from steel produced by closely controlled steelmaking practices and are subject to mill testing and inspection designed to ensure internal soundness, uniformity of chemical composition, and freedom from detrimental surface imperfections. Proper control of hardness and microstructure by heat treatment and cold working is important for cold forming. Most cold-heading quality alloy steels are low- and medium-carbon grades. Typical low-carbon alloy steel parts, made by cold heading, include fasteners (cap screws, bolts, eyebolts), studs, anchor pins, and rollers for bearings. Examples of medium-carbon alloy steel cold-headed parts are bolts, studs, and hexagon-headed cap screws. Special cold-heading quality applies to cold-finished alloy steel bars for applications involving severe cold plastic deformation when slight surface imperfections may cause splitting of a part. Bars of this quality are produced by closely controlled steelmaking practices to provide uniform chemical composition and internal soundness. Also, special processing (such as grinding) is applied at intermediate stages to remove detrimental surface imperfections. Proper control of hardness and microstructure by heat treatment and cold working is important for cold forming. Typical applications of alloy steel bars of this quality are front suspension studs, socket screws, and some valves. Axle shaft quality applies to cold-finished alloy steel bars intended for the manufacture of automotive or truck-type, power-driven axle shafts, which by their design or method of manufacture are either not machined all over or undergo less than the recommended amount of stock removal for proper cleanup of normal surface imperfections. Axle shaft quality bars require special rolling practices, special billet and bar conditioning, and selective inspection techniques. Ball and roller bearing quality and bearing quality apply to cold-finished alloy steel bars used for the manufacture of antifriction bearings. Such bars are usually produced from alloy steels of the AISI-SAE standard alloy carburizing grades and the AISI-SAE high-carbon chromium series. These steels can be produced in accordance with ASTM A 534, A 295, and A 485. Bearing quality steels are subjected to restricted melting and special teeming, heating, rolling, cooling, and conditioning practices


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to meet rigid quality requirements. The steelmaking operations may include vacuum treatment. The foregoing requirements include thorough examination for internal imperfections by one or more of the following methods: macroetch testing, microscopic or ultrasonic examination for nonmetallic inclusions, and fracture testing. Aircraft quality and magnaflux quality apply to cold-finished alloy steel bars for important or highly stressed parts of aircraft and for other similar or corresponding purposes involving additional stringent requirements, such as magnetic particle inspection, additional discard, macroetch tests, and hardenability control. The meet these requirements, exacting steelmaking, rolling, and testing practices must be employed. These practices are designed to minimize detrimental inclusions and porosity. Phosphorus and sulfur are usually limited to 0.025% maximum. There are many aircraft parts and many parts for missiles and other rockets that require aircraft quality steel. The magnetic particle testing requirements given in AMS 2301 are sometimes specified for such applications. Other Alloy Steel Qualities. The quality descriptors listed below apply to cold-finished alloy steel bars intended for rifles, guns, shell, shot, and similar applications. They may have requirements for amount of discard, macroetch testing, surface requirements, or magnetic particle testing as indicated in the product specifications: • • • • • •

Armor-piercing (AP) shot quality AP shot magnaflux quality Gun quality Rifle barrel quality Shell quality Shell magnaflux quality

Mechanical Properties A major difference between machined and cold-drawn round bars is the improvement in tensile and yield strengths that results from the cold work of drawing. Cold work also changes the shape of the stress-strain diagram, as shown in Fig. 1 . Within the range of commercial drafts, cold work markedly affects certain mechanical properties (Fig. 2 ). The variations in percentage of reduction of cross section for bars drawn with normal commercial drafts of 0.8 and 1.6 mm (1=32and 1=16in.) and with heavy drafts of 3.2 and 4.8 mm (1=8and 3=16in.) are shown in Fig. 3 . Normal reductions seldom exceed 20% and are usually less than 12%. According to Fig. 2 , the more pronounced changes in significant tensile properties occur within this range of reductions (up to about 15%). Fig. 1 Effect of cold work on the tensile stress-strain curve for low-carbon steel bars


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Fig. 2 Effect of cold drawing on the tensile properties of steel bars. Data are for bars up to 25 mm (1 in.) in cross section having a tensile strength of 690 MPa (100 ksi) or less before cold drawing.

Fig. 3 Effect of draft on reduction of cross section of steel bars


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The minimum mechanical properties of several cold-drawn carbon steel bars in a range of sizes are presented in Table 7 . In addition, the effects of both low- and high-temperature stress relief on the as-cold-drawn mechanical properties are noted. The mechanical property ranges and average values for one of the steels listed in Table 7 (1137 resulfurized steel) are presented in Fig. 4 , which also shows the advantage in strength of cold-drawn over hot-rolled material. Table 7 Estimated minimum mechanical properties of cold-drawn carbon steel rounds, squares, and hexagons Estiminated minimum mechanical properties for sizes under 16 mm (5=8in.) can be obtained from individual producers. The data in this table are not applicable to turned and polished or turned and ground bars, which have mechanical properties corresponding to those of hot-rolled steel bars of the same size and grade. The size of a square or hexagon is the distance between opposite sides. Cold drawn followed by Cold drawn followed by As cold-drawn low-temperature stress relief high-temperature stress relief Steel designation and size range

mm

in.

Strength

Strength

Elo nga tion in 50 mm Ha (2 rdn in.), ess, MP MP HB a ksi a ksi %

Elo nga tion in 50 mm Ha (2 rdn in.), ess, MP MP HB a ksi a ksi %

Strength

Elo nga tion in 50 mm (2 in.), % MP MP a ksi a ksi % Tensile

Yield

Tensile

Re duc tion in are a, %

Yield

Tensile


Yield


Ha rdn ess, HB

1018, 1025 16−22 inclusive

5

=8−7=8 inclusive

483 70 413 60

18

40

43

...

...

...

...

...

...

... 448 65 310 45

20

45 131

22−22 inclusive


448 65 379 55

16

40 131

...

...

...

...

...

...

... 414 60 310 45

20

45 121

32−51 inclusive

>11=4−2 inclusive

414 60 345 50

15

35 121

...

...

...

...

...

...

... 379 55 310 45

16

40 111

51−76 inclusive

>2−3 inclusive

379 55 310 45

15

35 111

...

...

...

...

...

...

... 345 50 276 40

15

40 101

1117, 1118 16−22 inclusive

5

=8−7=8 inclusive

517 75 448 65

15

40 149 552 80 483 70

15

40 163 483 70 345 50

18

45 143

22−32 inclusive


483 70 414 60

15

40 143 517 75 448 65

15

40 149 448 65 345 50

16

45 131

32−51

>11=4−2

448 65 379 55

13

35 131 483 70 414 60

13

35 143 414 60 345 50

15

40 121


Page 423



01 Sep 2005

inclusive

inclusive

51−76 inclusive

>2−3 inclusive

414 60 345 50

12

30 121 448 65 379 55

12

35 131 379 55 310 45

15

40 111

1035 16−22 inclusive

5

=8−7=8 inclusive

586 85 517 75

13

35 170 621 90 552 80

13

35 179 552 80 414 60

16

45 163

22−23 inclusive


552 80 483 70

12

35 163 586 85 517 75

12

35 170 517 75 414 60

15

45 149

32−51 inclusive

>11=4−2 inclusive

517 75 448 65

12

35 149 552 80 483 70

12

35 163 483 70 414 60

15

40 143

51−76 inclusive

>2−3 inclusive

483 70 414 60

10

30 143 517 75 448 65

10

30 149 448 65 379 55

12

35 131

1040, 1140 16−22 inclusive

5

=8−7=8 inclusive

621 91 552 80

12

35 179 655 95 586 85

12

35 187 586 85 448 65

15

45 170

22−32 inclusive


586 85 517 75

12

35 170 621 90 552 80

12

35 179 552 80 448 65

15

45 163

32−51 inclusive

>11=4−2 inclusive

552 80 483 70

10

30 163 586 85 517 75

10

30 170 517 75 414 60

15

40 149

51−76 inclusive

>2−3 inclusive

517 75 448 65

10

30 149 552 80 483 70

10

30 163 483 70 379 55

12

35 143

1045, 1145 16−22 inclusive

5

=8−7=8 inclusive

655 95 586 85

12

35 187 689 100 621 90

12

35 197 621 90 483 70

15

45 179

22−32 inclusive


621 90 552 80

11

30 179 655 95 586 85

11

30 187 586 85 483 70

15

45 170

32−51 inclusive

>11=4−2 inclusive

586 85 517 75

10

30 170 621 90 552 80

10

30 179 552 80 448 65

15

40 163

51−76 inclusive

>2−3 inclusive

552 80 483 70

10

30 163 586 85 517 75

10

25 170 517 75 414 60

12

35 149

1050, 1137, 1151 16−22 inclusive

5

=8−7=8 inclusive

689 100 621 90

11

35 197 724 105 655 95

11

35 212 655 95 517 75

15

45 187

22−32 inclusive


655 95 586 85

11

30 187 689 100 621 90

11

30 197 621 90 517 75

15

40 179

32−51 inclusive

>11=4−2 inclusive

621 90 552 80

10

30 179 655 95 586 85

10

30 187 586 85 483 70

15

40 170

51−76 inclusive

>2−3 inclusive

586 85 517 75

10

30 170 621 90 552 80

10

25 179 552 80 448 65

12

35 163


5

=8−7=8 inclusive

724 105 655 95

11

30 212 758 110 689 100 11

30 223 689 100 552 80

15

40 197

22−32 inclusive


689 100 621 90

10

30 197 724 105 655 95

10

30 212 655 95 552 80

15

40 187

32−51 inclusive

>11=4−2 inclusive

655 95 586 85

10

30 187 689 100 621 90

10

25 197 621 90 517 75

15

40 179

51−76 inclusive

>2−3 inclusive

621 90 552 80

10

20 179 655 95 581 85

10

20 187 586 85 483 70

12

30 170


5

=8−7=8 inclusive

758 110 689 100 10

30 223 793 115 724 105 10

30 229 724 105 586 85

15

40 212

22−32 inclusive


724 105 655 95

10

30 212 758 110 689 100 10

30 223 689 100 586 85

15

40 197

32−51

>11=4−2

689 100 621 90

10

25 197 724 105 655 95

25 212 655 95 552 80

15

35 187


10

Page 424


inclusive


01 Sep 2005

inclusive

51−76 >2−3 inclusive inclusive Source: Ref 2

655 95 586 85

10

20 187 689 100 621 90

10

20 197 621 90 517 75

12

30 179

Fig. 4 Mechanical properties of hot-rolled and cold-drawn 1137 bars

Measurements of the changes in mechanical properties of three specific carbon steels (1016, 1040, and 1060) and one alloy steel (8630) as a result of cold-drawing are shown in Fig. 5 , 6 , 7 , and 8 . Some of these data pertain to large reductions, well beyond commercial ranges. Data plotted as solid lines in Fig. 5 , 6 , 7 , and 8 are for bars that were cold drawn and artificially aged, but not stress relieved. After drawing, these bars were aged for 4 h at 100 °C (212 °F) to simulate the natural aging resulting from several months of storage at room temperature. Data plotted as dashed lines in Fig. 5 , 6 , 7 , and 8 are for cold-drawn and stress-relieved bars; carbon steels were stress-relieved for 2 h at 480 °C (900 °F), and 8630 steel for 2 h at 540 °C (1000 °F), after cold drawing. Fig. 5 Effects of cold drawing and of cold drawing and stress relieving on mechanical properties of 1016 steel bars. Dashed curves represent cold-drawn material; solid curves, material cold drawn and stress relieved. All bars were from a single heat. The bars were hot reduced to a diameter of 51 mm (2 in.) by conventional practice, then normalized and cold drawn to the reductions indicated. Note that the larger reductions are well beyond commercial ranges. Test specimens were taken from half-radius positions in the bars. The bars that were stress relieved after cold drawing were treated as follows: carbon steel bars, 2 h at 480 °C (900 °F); 8630 bars, 2 h at 540 °C (1000 °F). All bars were aged for 4 h at 100 °C (212 °F) after cold drawing, or after cold drawing and stress relieving, to simulate the condition of steel after several months of storage at room temperature. Source: Ref 5


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01 Sep 2005

Fig. 6 Effects of cold drawing and of cold drawing and stress relieving on mechanical properties of 1040 steel bars. Dashed curves represent cold-drawn material; solid curves, material cold drawn and stress relieved. All bars were from a single heat. The bars were hot reduced to a diameter of 51 mm (2 in.) by conventional practice, then normalized and cold drawn to the reductions indicated. Note that the larger reductions are well beyond commercial ranges. Test specimens were taken from half-radius positions in the bars. The bars that were stress relieved after cold drawing were treated as follows: carbon steel bars, 2 h at 480 °C (900 °F); 8630 bars, 2 h at 540 °C (1000 °F). All bars were aged for 4 h at 100 °C(212 °F) after cold drawing, or after cold drawing and stress relieving, to simulate the condition of steel after several months of storage at room temperature. Source: Ref 5


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01 Sep 2005

Fig. 7 Effects of cold drawing and of cold drawing and stress relieving on mechanical properties of 1060 steel bars. Dashed curves represent cold-drawn material; solid curves, material cold drawn and stress relieved. All bars were from a single heat. The bars were hot reduced to a diameter of 51 mm (2 in.) by conventional practice, then normalized and cold drawn to the reductions indicated. Note that the larger reductions are well beyond commercial ranges. Test specimens were taken from half-radius positions in the bars. The bars that were stress relieved after cold drawing were treated as follows: carbon steel bars, 2 h at 480 °C (900 °F); 8630 bars 2 h at 540 °C (1000 °F). All bars were aged for 4 h at 100 °C (212 °F) after cold drawing, or after cold drawing and stress relieving, to simulate the condition of steel after several months of storage at room temperature. Source: Ref 5


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01 Sep 2005

Fig. 8 Effects of cold drawing and of cold drawing and stress relieving on mechanical properties of 8630 steel bars. Dashed curves represent cold-drawn material; solid curves, material cold drawn and stress relieved. All bars were from a single heat. The bars were hot reduced to a diameter of 51 mm (2 in.) by conventional practice, then normalized and cold drawn to the reductions indicated. Note that the larger reductions are well beyond commercial ranges. Test specimens were taken from half-radius positions in the bars. The bars that were stress relieved after cold drawing were treated as follows: carbon steel bars, 2 h at 480 °C (900 °F); 8630 bars 2 h at 540 °C (1000 °F). All bars were aged for 4 h at 100 °C (212 °F) after cold drawing, or after cold drawing and stress relieving, to simulate the condition of steel after several months of storage at room temperature. Source: Ref 5


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01 Sep 2005

Tensile and Yield Strengths. The data in Fig. 2 , 5 , 6 , 7 , and 8 indicate that as cold work increases up to about 15%, yield strength increases at a greater rate than tensile strength. The greatest improvement in strength results from the first 5% of reduction. Stress relieving modifies this pattern appreciably. They yield/tensile ratio is markedly affected by cold drawing. In this condition (cold drawn, not stress relieved; see Fig. 5 , 6 , 7 , and 8 ), the data for yield/tensile ratio indicate a somewhat erratic behavior. However, the ratio follows a consistent upward trend with increased cold work and subsequent stress relief. The hardness increases with increased cold work and, in most cases, is affected by stress relieving. There is considerable scatter in the relations between hardness and tensile strength and hardness and yield strength, as indicated by the data in Fig. 9 for 41 heats of cold-drawn and stress-relieved 1144 steel. However, there is a relationship between hardness and tensile strength or hardness and yield strength, because published tables allow approximation of the hardness or tensile strength (or yield strength) when one of the other values is known. Fig. 9 Mechanical properties of 1144 steel bars cold drawn and stress relieved at 565 °C (1050 °F). Range of composition for 41 heats was 0.41 to 0.52% C, 1.33 to 1.68% Mn, and 0.220 to 0.336% S.


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01 Sep 2005

Impact Properties. Available data are limited on the effect of cold work on notched-bar impact properties. The results of one of the more important studies are included in Fig. 5 , 6 , 7 , and 8 , which show the effect of cold work over a wide range of drafts on three carbon steels with increasing carbon contents and the effect of cold work on 8630 alloy steel. Within the range of commercial drafts, energy absorbed (breaking strength) falls rapidly for the 1016 steel and less rapidly for 8630 steel. At any level of cold work, energy absorbed decreases with increased carbon content. In the stress-relieved condition, the fracture transition temperature generally rises with increasing amounts of cold work up to 20 to 30% reduction. Beyond this commercial range of reductions, the transition temperature falls. For 1016 steel, extremely heavy drafts lower the transition temperature to below that of the original hot-rolled material. Increasing carbon content raises the transition temperature.

Residual Stresses The stress pattern produced by cold drawing depends on the amount of reduction and the shape of the die, as well as the microstructure, hardness, and grade of steel. Figure 10 illustrates the effect of reduction in area on the magnitude and distribution of stresses in bars of 1050 steel reduced by the amounts shown. Cold drawing of the bars to 4.1% reduction resulted in surface compressive stresses, while increasing the amount of cold drawing to 12.3% reduction resulted in a change of the surface stresses from compressive to tensile. The variation in longitudinal stress over a much wider range of reduction values is shown in Fig. 11 for steel wire (the effect is qualitatively similar for bars). The greatest effect on the residual stress is caused by the first 10% reduction. The effect of a very light draft is to produce compressive stress at the surface, which rapidly changes to tensile stress with a relatively small increase in reduction. Fig. 10 Effect of increasing reduction on the residual stress patterns in cold-drawn bars of 1050 steel. (a) 4.1% reduction in area. (b) 6.2% reduction in area. (c) 8.3% reduction in area. (d) 12.3% reduction in area. Source: Ref 6


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01 Sep 2005

Fig. 11 Effect of increasing single-draft reduction on residual longitudinal stress at the surface of drawn steel wire

Both straightening and stress relieving after cold drawing have significant effects on the residual stress pattern of the resulting product. Figure 12 shows the longitudinal, tangential, and radial stress patterns that result when a 43 mm (111=16in.) diam carbon steel bar is drawn to 38 mm (11=2in.), a reduction of 20%. These data indicate that the surface of the bar is in tension, the center is in compression, and both longitudinal and tangential stresses vary over wide ranges. Fig. 12 Residual stress patterns obtained in cold-drawn steel bars of 1045 steel. (a) As-drawn. (b) After rotary straightening. (c) After stress relieving. Bars were cold drawn 20% from 43 to 38 mm (111=16to 11=2in.).


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01 Sep 2005

Straightening in a skewed-rolls (Medart) machine significantly reduces residual stress, particularly at the surface, as shown in Fig. 12 (b). It is of interest to compare the longitudinal stress curve shown in Fig. 12 (b) with that in Fig. 12 (a). Figure 12 (c) shows the effect of two stress-relieving temperatures (425 and 540 °C, or 800 and 1000 °F) on residual longitudinal stress. Stress relieving at these temperatures is only slightly more effective than straightening in reducing the residual stress level. This phenomenon may be accounted for by an analysis of the nature of the stresses that are developed in cold drawing. The stress applied in cold drawing is sufficient to deform the material both elastically and plastically. Because the initiation of plastic strain depends on the development of maximum elastic strain, the ratio of these two strains after release of the deforming stress may be highly variable. If the deformation caused by cold drawing were uniform across the section, as in pure stretching, the elastic stress would be released by the release of the deforming stress. Because the plastic strain is not uniform, as shown by the dishing of the ends of drawn bars, neither is the accompanying elastic strain. When the deforming stress is removed from such a system, the remaining nonuniform elastic-strain energy cannot be released completely, because the resistance of low-strain regions prevents the complete recovery of regions of high strain. A pattern of residual stress results from this unequal adjustment. Stress Relieving. The inevitable residual stresses in as-drawn bars can be relieved mechanically or thermally (Ref 7). Mechanical relief may take two forms. One involves the introduction of stresses of opposite sign, which can be accomplished by shot peening. A second approach is to plastically deform the material further, thus affording additional opportunity for the relief of non-uniform residual stresses. The data on rotary straightening in Fig. 12 (b) demonstrate this effect. The thermal stress relieving of cold-drawn bars⎯also known as strain drawing, strain annealing, strain relieving, preaging, and stabilizing⎯is probably the most widely used thermal treatment applied to cold-drawn bars. Its purpose is to modify the magnitude and distribution of residual stresses in the cold-finished bar and thus produce a product with the desired combination of mechanical properties for field service. Thermal stress relief temporarily reduces the strength level of the material (at the stress-relieving temperature) and enables the elastic-strain energy to find release in small but significant amounts of plastic deformation. After stress relieving, the maximum residual stress that can remain is equal to the yield strength of the material at the stress-relieving temperature. Temperatures up to about 650 °C (1200 °F) are commonly used for the stress relieving of cold-drawn bars. The upper limit for the stress-relief temperature for a particular cold-worked steel is the recrystallization or lower critical temperature of that steel, because if this temperature is exceeded, the strengthening effect of cold work is lost. The temperatures used in commercial practice frequently range from 370 to 480 °C (700 to 900 °F). When stress relieving is performed at relatively low temperatures (for example, 290 °C, or 550 °F), yield strength of most cold-drawn steels is increased. At higher temperatures, however, hardness, tensile strength, and yield strength are reduced, while elongation and reduction in area are increased. The choice of a specific time and temperature is dependent on chemical composition, cold-drawing practice, and the final properties required in the bar. The various categories of stress relief can be divided into three groups: • Group 1: Complete relief of all cold-working stresses • Group 2: Relief of cold-working stresses to a limited degree to increase ductility and stability in the material • Group 3: Relief of stresses in heavily drafted steels to develop high yield strength Group 1 treatment is conducted above 540 °C (1000 °F). It removes all residual stresses that otherwise would cause objectionable distortion in machining. Group 2 Treatment. With group 2 processing, lower temperatures in the range of 370 to 540 °C (700 to 1000 °F) are used, and the stresses are partially relieved to bring the mechanical properties within the limits of individual specifications. Applications falling into this class are those that may require ductility close to that of hot-rolled steel, along with good surface finish and close


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01 Sep 2005

control of dimensions and stability during machining. Group 3 stress relief is used for bars with heavy drafts. These drafts raise the tensile and yield strengths to high levels, but reduce elongation and reduction in area. Heating to 260 to 425 °C (500 to 800 °F) restores the ductility while retaining or increasing the strength and hardness imparted by the cold work. The effect of stress relieving at two temperatures on the residual stress pattern of 38 mm (1.5 in.) diam 0.45% C steel bars that have been cold drawn 20% is shown in Fig. 12 . The estimated minimum mechanical properties for cold-drawn carbon steel bars as-cold-drawn and as-cold-drawn followed by both a low- and a high-temperature stress-relieving treatment are given in Table 7 . The cumulative effects of cold drawing, straightening, and stress relieving on the yield and tensile strength of 1144 steel are shown in Fig. 13 . Fig. 13 Various production stages of 1144 steel. A, hot rolled; B, cold drawn; C, cold drawn and straightened; D, cold drawn, straightened, and strain relieved

Heat Treatment Heat treatment by quenching and tempering, followed by scale removal and then cold drawing, can also be used as a method of producing stronger cold-finished bars in those grades amenable to quench hardening. Heat treatment provides the required increase in strength, and cold drawing provides the size and finish, with a minimal increase in the mechanical properties obtained by quench hardening. Alternatively, quenched and tempered bars can be cold finished by turning and polishing. When bars are cold finished by turning and polishing, there is no increase in the mechanical properties obtained by quench hardening. For the cold drawing of quenched and tempered bars to be economically justifiable, the minimum strength level produced must be above that obtainable by conventional cold-drawn practices. The upper strength limit is not clearly defined, but for most applications it is the upper limit of machinability. The cold drawing of quenched and tempered bars is applicable to both carbon and alloy steels; however, for the process to be economically justifiable, alloy steel is used only in those sizes above which carbon steel will not respond satisfactorily to liquid quenching. Typically, quenched and tempered product offers superior ductility and heat-resistant properties. Other heat treatments⎯principally normalizing, full annealing, spheroidizing, and thermal stress relieving⎯can be applied to suitable grades of hot-rolled steel before or after cold drawing or turning and polishing as required by the end product. Control of microstructure is frequently important, a good example being annealing for machinability. A controlled rate of


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01 Sep 2005

continuous cooling through the pearlite transformation range (so-called cycle annealing) is employed. Isothermal cycles are also used. In the cycle-annealing process, the rate of cooling of the furnace charge is adjusted so that the time required to traverse the pearlite temperature interval is sufficient to allow completion of that transformation. By regulating the dwell in the transformation temperature range, the carbide distribution in the product can be varied from partly spheroidal to fully pearlitic, and the pearlite from coarse to fine. In this manner, the optimum machining structure can be obtained for the grade and the machining practice being used. Spheroidize annealing thermal treatment is given to cold-finished bars that are to be used for severe cold-forming operations. The aim of this treatment is to develop a microstructure consisting of globular carbides in a ferrite matrix. The rate of spheroidizing depends to some degree on the original microstructures. Prior cold work also increases the rate of spheroidizing, particularly for subcritical spheroidizing treatments. The spheroidized structure is desirable when minimum hardness and maximum ductility are important. Low-carbon steels are seldom annealed for machining because, in the annealed condition, they are very soft and gummy, which tends to produce long, stringy chips that cause handling problems at the machine tool and contribute to a rough surface finish on the machined part. When such steels are spheroidized, it is usually to permit severe cold deformation. Carbon Restoration. During the hot-working operations involved in the production of bar products⎯the reduction of cast ingots, blooms, or billets and subsequent conversion in bar mills⎯decarburization of the bar surface takes place because of exposure to ambient oxygen at high temperatures throughout these operations. A specialized variant of full annealing, called carbon restoration or carbon correction, is utilized to compensate for the loss of carbon due to decarburization. Carbon restoration for alloy steels is limited because vanadium carbide and molybdenum are not recovered. By heating the descaled hot-rolled bars to approximately 870 to 925 °C (1600 to 1700 °F) in a controlled atmosphere, it is possible to restore surface carbon to the required level. A modern controlled-atmosphere furnace is used for this purpose. Methane or other light hydrocarbons are burned with a controlled amount of air in an endothermic generator to produce a gas with a mixed ratio of CO to CO2. By controlling the CO/CO2 ratio of the endothermic gas, an atmosphere can be generated that will be in equilibrium with the carbon content of the steel to be treated. Low-ratio gas is in equilibrium with lower-carbon steels, and high-ratio gas is in equilibrium with higher content steels. The actual ratio used depends on the type of anneal and the grade of steel to be annealed. This ratio must be closely controlled, or the atmosphere will become decarburizing or carburizing to the steel. Modern instruments, such as oxygen probes, are available to maintain this close control. After carbon restoration, bars are cold drawn. Material processed in this manner is useful when parts must have full hardness on the cold-drawn (unmachined) surface after heat treatment. Many induction-hardened parts make use of a carbon-restored material as a means of eliminating machining.

Machinability Cold drawing significantly improves the machinability of the steels discussed in this article. The increase in hardness due to cold work causes the chips formed by a cutting tool to tear away from the workpiece more readily, and to be harder and more brittle, so that they break up easily and are less likely to build up on the tool edge. Deformation extends a shorter distance above the edge of the tool, giving a sharper cleavage at that point. These factors contribute to improvements in power consumption, tool wear, and surface finish. They result from the addition of the major contributors to improvements in machinability: phosphorus, sulfur, nitrogen (diatomic), lead, bismuth, tellurium, selenium, calcium, and so on, in various combinations. In addition, the accuracy of size and section of cold-finished bars minimizes collet troubles and requires less surface removal to obtain concentricity. The freedom from scale on the cold-finished bar also improves tool life and may permit the surface of the bar to be used as the finished surface of the completed part. Cold drawing generally improves the machinability of low-carbon steels because the high ductility of these materials in the hot-rolled condition can be lowered considerably without raising strength excessively. In contrast, a steel such as 1144, which is inherently low in ductility because of its higher carbon content, shows little improvement in machining after cold drawing. The increased hardness that results from cold drawing can be deleterious to the machinability of the higher-carbon steels; it may be helpful to stress relieve after cold drawing to reduce hardness. Another approach to maximum machinability with the higher-carbon grades is to anneal before cold drawing. This puts the carbide in a form that is less abrasive to the cutting tool. Lamellar anneal and spheroidize annealing are used depending on carbon level, machinability requirements, and heat treat response requirements. The trade-off values must be decided for each individual application. Compared with hot-rolled steel, uniformity of hardness and structure are improved. One of the conventional indexes of machinability is the ratio of tool life to that encountered with 1212 cold-drawn bars. The average machinability ratings for cold-drawn carbon steel bars, nonresulfurized and resulfurized carbon steel bars, and alloy steel bars, based on a value of 100% for 1212 bars, are given in Tables 8 , 9 , 10 , and 11 . The relative machinability data listed in Tables 8 , 9 , 10 , and 11 represent results obtained from experimental studies and actual shop production information on the general run of parts. Any extraordinary features of the part to be machined or physical conditions of the steel should be taken into consideration, and speeds and feeds altered accordingly. In addition, machinability is influenced by various metallurgical factors, such as degree of cold reduction, mechanical properties, grain size, and microstructure. Therefore, the data in Tables 8 , 9 , 10 , and 11 are presented only as a starting point from which proper speeds and feeds for specific parts can be determined. Further discussion of the machinability of cold-drawn steel is included in the article "Machinability of Steels" in this Volume. Table 8 Machinability ratings and recommended feeds and speeds for cold-drawn carbon steel bars


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01 Sep 2005

All cutting speeds and feeds based on cutting with high-speed steel tools. For cutting conditions for other machining operations, see the article "Machinability of Steels" in this Volume. Form turning Single-point turning Drilling Machin Steel designation AISI or SAE 12L14

1213, 1215

1119, 1212

1211

1117, 1118

1144, annealed

ability rating( a), %

Width of cut, mm (in.)

Speed, m/min (sfm)

Feed, mm/rev (in./rev)

Depth of cut, mm (in.)

Speed, m/min (sfm)


Size of hole, mm (in.)

Speed, m/min (sfm)


170

12.7 (0.500)

85.3 (280)

0.089 (0.0035)

3.18 (0.125)

85.3 (280)

0.236 (0.0093)

6.35 (0.250)

50.3 (165)

0.160 (0.0063)

25.4 (1.000)

79.2 (260)

0.076 (0.0030)

6.35 (0.250)

79.2 (260)

0.224 (0.0088)

12.7 (0.500)

50.3 (165)

0.175 (0.0069)

38.1 (1.500)

79.2 (260)

0.074 (0.0029)

9.53 (0.375)

77.7 (255)

0.180 (0.0071)

19.05 (0.750)

54.9 (180)

0.206 (0.0081)

50.8 (2.000)

76.2 (250)

0.053 (0.0021)

12.7 (0.500)

76.2 (250)

0.152 (0.0060)

25.4 (1.000)

54.9 (180)

0.234 (0.0092)

63.5 (2.500)

71.6 (235)

0.043 (0.0017)

... ...

... ...

... ...

31.75 (1.250)

56.4 (185)

0.267 (0.0105)

12.7 (0.500)

68.6 (225)

0.076 (0.0030)

3.18 (0.125)

68.6 (225)

0.216 (0.0085)

6.35 (0.250)

38.1 (125)

0.137 (0.0054)

25.4 (1.000)

64.0 (210)

0.064 (0.0025)

6.35 (0.250)

64.0 (210)

0.203 (0.0080)

12.7 (0.500)

38.1 (125)

0.152 (0.0060)

38.1 (1.500)

64.0 (210)

0.064 (0.0025)

9.53 (0.375)

62.5 (205)

0.165 (0.0065

19.05 (0.750)

42.7 (140)

0.178 (0.0070)

50.8 (2.000)

62.5 (205)

0.046 (0.0018)

12.7 (0.500)

61.0 (200)

0.140 (0.0055

25.4 (1.000)

42.7 (140)

0.203 (0.0080)

63.5 (2.500)

61.0 (200)

0.038 (0.0015)

... ...

... ...

... ...

31.75 (1.250)

44.2 (145)

0.229 (0.0090)

12.7 (0.500)

50.3 (165)

0.064 (0.0025)

3.18 (0.125)

50.3 (165)

0.178 (0.0070)

6.35 (0.250)

32.0 (105)

0.114 (0.0045)

25.4 (1.000)

48.8 (160)

0.051 (0.0020)

6.35 (0.250)

48.8 (160)

0.165 (0.0065

12.7 (0.500)

32.0 (105)

0.127 (0.0050)

38.1 (1.500)

48.8 (160)

0.046 (0.0018)

9.53 (0.375)

47.2 (155)

0.140 (0.0055)

19.05 (0.750)

35.0 (115)

0.152 (0.0060)

50.8 (2.000)

47.2 (155)

0.038 (0.0015)

12.7 (0.500)

45.7 (150)

0.114 (0.0045)

25.4 (1.000)

35.0 (115)

0.178 (0.0070)

63.5 (2.500)

45.7 (150)

0.030 (0.0012)

... ...

... ...

... ...

31.75 (1.250)

36.6 (120)

0.203 (0.0080)

12.7 (0.500)

47.2 (155)

0.058 (0.0023)

3.18 (0.125)

47.2 (155)

0.168 (0.0066)

6.35 (0.250)

30.2 (99)

0.107 (0.0042)

25.4 (1.000)

45.7 (150)

0.048 (0.0019)

6.35 (0.250)

45.7 (150)

0.155 (0.0061)

12.7 (0.500)

30.2 (99)

0.119 (0.0047)

38.1 (1.500)

45.7 (150)

0.043 (0.0017)

9.53 (0.375)

44.5 (146)

0.132 (0.0052)

19.05 (0.750)

32.9 (108)

0.142 (0.0056)

50.8 (2.000)

44.5 (146)

0.036 (0.0014)

12.7 (0.500)

43.0 (141)

0.107 (0.0042)

25.4 (1.000)

32.9 (108)

0.168 (0.0066)

63.5 (2.500)

43.0 (141)

0.028 (0.0011)

... ...

... ...

... ...

31.75 (1.250)

34.4 (113)

0.193 (0.0076)

12.7 (0.500)

45.7 (150)

0.056 (0.0022)

3.18 (0.125)

45.7 (150)

0.163 (0.0064)

6.35 (0.250)

29.0 (95)

0.104 (0.0041)

25.4 (1.000)

44.2 (145)

0.046 (0.0018)

6.35 (0.250)

44.2 (145)

0.150 (0.0059)

12.7 (0.500)

29.0 (95)

0.114 (0.0045)

38.1 (1.500)

44.2 (145)

0.041 (0.0016)

9.53 (0.375)

43.0 (141)

0.127 (0.0050)

19.05 (0.750)

32.0 (105)

0.140 (0.0055)

50.8 (2.000)

43.0 (141)

0.036 (0.0014)

12.7 (0.500)

41.5 (136)

0.104 (0.0041)

25.4 (1.000)

32.0 (105)

0.163 (0.0064)

63.5 (2.500)

41.5 (136)

0.028 (0.0011)

... ...

... ...

... ...

31.75 (1.250)

36.3 (119)

0.185 (0.0073)

12.7

42.7

0.053

3.18

42.7

0.150

6.35

27.1 (89)

0.102

136

100

94

91

85


Page 435


1141, annealed

1016, 1018, 1022

1144

1020, 1137, 1045, annealed

1035, 1141, 1050, annealed

81

78

76

72

70


01 Sep 2005

(0.500)

(140)

(0.0021)

(0.125)

(140)

(0.0059)

(0.250)

25.4 (1.000)

41.5 (136)

0.043 (0.0017)

6.35 (0.250)

41.5 (136)

0.140 (0.0055)

12.7 (0.500)

27.1 (89)

0.114 (0.0045)

38.1 (1.500)

41.5 (136)

0.038 (0.0015)

9.53 (0.375)

40.2 (132)

0.119 (0.0047)

19.05 (0.750)

29.9 (98)

0.140 (0.0055)

50.8 (2.000)

40.2 (132)

0.033 (0.0013)

12.7 (0.500)

38.7 (127)

0.102 (0.0040)

25.4 (1.000)

29.9 (98)

0.163 (0.0064)

63.5 (2.500)

38.7 (127)

0.25 (0.0010)

... ...

... ...

... ...

31.75 (1.250)

31.1 (102)

0.179 (0.0070)

12.7 (0.500)

41.1 (135)

0.051 (0.0020)

3.18 (0.125)

41.1 (135)

0.145 (0.0057)

6.35 (0.250)

26.2 (86)

0.102 (0.0040)

25.4 (1.000)

39.6 (130)

0.043 (0.0017)

6.35 (0.250)

39.6 (130)

0.135 (0.0053)

12.7 (0.500)

26.2 (86)

0.114 (0.0045)

38.1 (1.500)

39.6 (130)

0.038 (0.0015)

9.53 (0.375)

38.7 (127)

0.114 (0.0045)

19.05 (0.750)

28.6 (94)

0.137 (0.0054)

50.8 (2.000)

38.7 (127)

0.030 (0.0012)

12.7 (0.500)

37.2 (122)

0.094 (0.0037)

25.4 (1.000)

28.6 (94)

0.160 (0.0063)

63.5 (2.500)

37.2 (122)

0.025 (0.0010)

... ...

... ...

... ...

31.75 (1.200)

29.9 (98)

0.183 (0.0072)

12.7 (0.500)

39.6 (130)

0.048 (0.0019)

3.18 (0.125)

39.6 (130)

0.140 (0.0055)

6.35 (0.250)

25.0 (82)

0.096 (0.0038)

25.4 (1.000)

38.1 (125)

0.041 (0.0016)

6.35 (0.250)

38.1 (125)

0.130 (0.0051)

12.7 (0.500)

25.0 (82)

0.109 (0.0043)

38.1 (1.500)

38.1 (125)

0.036 (0.0014)

9.53 (0.375)

36.9 (121)

0.110 (0.0043)

19.05 (0.750)

27.4 (90)

0.132 (0.0052)

50.8 (2.000)

36.9 (121)

0.030 (0.0012)

12.7 (0.500)

35.7 (117)

0.090 (0.0035)

25.4 (1.000)

27.4 (90)

0.152 (0.0060)

63.5 (2.500)

35.7 (117)

0.023 (0.0009)

... ...

... ...

... ...

31.75 (1.250)

28.7 (94)

0.173 (0.0068)

12.7 (0.500)

38.1 (125)

0.048 (0.0019)

3.18 (0.125)

38.1 (125)

0.132 (0.0052)

6.35 (0.250)

24.1 (79)

0.094 (0.0037)

25.4 (1.000)

36.9 (121)

0.038 (0.0015)

6.35 (0.250)

36.9 (121)

0.124 (0.0049)

12.7 (0.500)

24.1 (79)

0.107 (0.0042)

38.1 (1.500)

36.9 (121)

0.036 (0.0014)

9.53 (0.375)

35.7 (117)

0.104 (0.0041)

19.05 (0.750)

26.5 (87)

0.127 (0.0050)

50.8 (2.000)

35.7 (117)

0.028 (0.0011)

12.7 (0.500)

34.4 (113)

0.086 (0.0034)

25.4 (1.000)

26.5 (87)

0.147 (0.0058)

63.5 (2.500)

34.4 (113)

0.023 (0.0009)

... ...

... ...

... ...

31.75 (1.250)

27.7 (91)

0.168 (0.0066)

12.7 (0.500)

36.6 (120)

0.046 (0.0018)

3.18 (0.125)

36.6 (120)

0.127 (0.0050)

6.35 (0.250)

23.2 (76)

0.089 (0.0035)

25.4 (1.000)

35.1 (115)

0.036 (0.0014)

6.35 (0.250)

35.1 (115)

0.119 (0.0047)

12.7 (0.500)

23.2 (76)

0.102 (0.0040)

38.1 (1.500)

35.1 (115)

0.033 (0.0013)

9.53 (0.375)

34.1 (112)

0.102 (0.0040)

19.05 (0.750)

25.3 (83)

0.119 (0.0047)

50.8 (2.000)

34.1 (112)

0.028 (0.0011)

12.7 (0.500)

32.9 (108)

0.081 (0.0032)

25.4 (1.000)

25.3 (83)

0.140 (0.0055)

63.5 (2.500)

32.9 (108)

0.023 (0.0009)

... ...

... ...

... ...

31.75 (1.250)

26.2 (86)

0.163 (0.0064)

12.7 (0.500)

35.1 (115)

0.043 (0.0017)

3.18 (0.125)

35.1 (115)

0.124 (0.0049)

6.35 (0.250)

22.3 (73)

0.086 (0.0034)

25.4 (1.000)

34.1 (112)

0.036 (0.0014)

6.35 (0.250)

34.1 (112)

0.114 (0.0045)

12.7 (0.500)

22.3 (73) 0.97 (0.0038)

38.1 (1.500)

34.1 (112)

0.033 (0.0013)

9.53 (0.375)

32.9 (108)

0.097 (0.0038)

19.05 (0.750)

24.4 (80)

0.114 (0.0045)

50.8 (2.000)

32.9 (108)

0.028 (0.0011)

12.7 (0.500)

32.0 (105)

0.079 (0.0031)

25.4 (1.000)

24.4 (80)

0.135 (0.0053)


(0.0040)

Page 436


1040

1045

1050

64

57

54


01 Sep 2005

63.5 (2.500)

32.0 (105)

0.020 (0.0008)

... ...

... ...

... ...

31.75 (1.250)

25.6 (84)

0.157 (0.0062)

12.7 (0.500)

32.0 (105)

0.038 (0.0015)

3.18 (0.125)

32.0 (105)

0.112 (0.0044)

6.35 (0.250)

20.4 (67)

0.081 (0.0032)

25.4 (1.000)

30.8 (101)

0.030 (0.0012)

6.35 (0.250)

30.8 (101)

0.104 (0.0041)

12.7 (0.500)

20.4 (67)

0.089 (0.0035)

38.1 (1.500)

30.8 (101)

0.038 (0.0011)

9.53 (0.375)

29.9 (98)

0.086 (0.0034)

19.05 (0.750)

22.3 (73)

0.107 (0.0042)

50.8 (2.000)

29.9 (98)

0.023 (0.0009)

12.7 (0.500)

29.0 (95)

0.071 (0.0028)

25.4 (1.000)

22.3 (73)

0.124 (0.0049)

63.5 (2.500)

29.0 (95)

0.018 (0.0007)

... ...

... ...

... ...

31.75 (1.250)

23.2 (76)

0.142 (0.0056)

12.7 (0.500)

29.0 (95)

0.036 (0.0014)

3.18 (0.125)

29.0 (95)

0.102 (0.0040)

6.35 (0.250)

18.3 (60)

0.071 (0.0028)

25.4 (1.000)

27.7 (91)

0.030 (0.0012)

6.35 (0.250)

27.7 (91)

0.094 (0.0037)

12.7 (0.500)

18.3 (60)

0.079 (0.0031)

38.1 (1.500)

27.7 (91)

0.025 (0.0010)

9.53 (0.375)

26.8 (88)

0.079 (0.0031)

19.05 (0.750)

19.8 (65)

0.094 (0.0037)

50.8 (2.000)

26.8 (88)

0.023 (0.0009)

12.7 (0.500)

25.9 (85)

0.066 (0.0026)

25.4(1.000) 19.8 (65)

0.112 (0.0044)

63.5 (2.500)

25.9 (85)

0.018 (0.0007)

... ...

... ...

... ...

31.75 (1.250)

20.7 (68)

0.127 (0.0050)

12.7 (0.500)

27.4 (90)

0.036 (0.0014)

3.18 (0.125)

27.4 (90)

0.097 (0.0038)

6.35 (0.250)

17.4 (57)

0.071 (0.0028)

25.4 (1.000)

26.5 (87)

0.028 (0.0011)

6.35 (0.250)

26.5 (87)

0.089 (0.0035)

12.7 (0.500)

17.4 (57)

0.079 (0.0031)

38.1 (1.500)

26.5 (87)

0.025 (0.0010)

9.53 (0.375)

25.6 (84)

0.076 (0.0030)

19.05 (0.750)

18.9 (62)

0.094 (0.0037)

50.8 (2.000)

25.6 (84)

0.020 (0.0008)

12.7 (0.500)

24.7 (81)

0.061 (0.0024)

25.4 (1.000)

18.9 (62)

0.112 (0.0044)

... ...

... ...

31.75 (1.250)

19.8 (65)

0.127 (0.0050)

63.5 24.7 (81) 0.018 ... ... (2.500) (0.0007) (a) Based on a machinability rating of 100% for 1212 steel. Source: Ref 2

Table 9 Estimated mechanical properties and machinability ratings of nonresulfurized carbon steel bars Estimated minimum values Steel designation SAE and/or AISI No.

UNS No.

Tensile strength Yield strength Type of processing (a)

MPa

ksi

MPa

ksi

Elongation in 50 mm Reduction in (2 in.), % area, %

Hardness, HB

Average machinab ility rating(b)

Manganese 1.00% maximum 1006 1008 1010 1012

G10060 G10080 G10100 G10120

Hot rolled

300

43

170

24

30

55

86

Cold drawn

330

48

280

41

20

45

95

Hot rolled

303

44

170

24.5

30

55

86

Cold drawn

340

49

290

41.5

20

45

95

Hot rolled

320

47

180

26

28

50

95

Cold drawn

370

53

300

44

20

40

105

Hot rolled

330

48

180

26.5

28

50

95

Cold drawn

370

54

310

45

19

40

105

1015

G10150

Hot rolled

340

50

190

27.5

28

50

101

Cold drawn

390

56

320

47

18

40

111

1016

G10160

Hot rolled

380

55

210

30

25

50

111

Cold drawn

420

61

350

51

18

40

121

1017

G10170

Hot rolled

370

53

200

29

26

50

105


50 55 55 55 60 70

Page 437



01 Sep 2005

Cold drawn

410

59

340

49

18

40

116

32

25

50

116

1018

G10180

Hot rolled

400

58

220

Cold drawn

440

64

370

54

15

40

126

1019

G10190

Hot rolled

410

59

220

32.5

25

50

116

Cold drawn

460

66

380

55

15

40

131

1020

G10200

Hot rolled

380

55

210

30

25

50

111

Cold drawn

420

61

350

51

15

40

121

1021

G10210

Hot rolled

420

61

230

33

24

48

116

Cold drawn

470

68

390

57

15

40

131

1022

G10220

Hot rolled

430

62

230

34

23

47

151

Cold drawn

480

69

400

58

15

40

137

1023

G10230

Hot rolled

370

56

210

31

25

50

111

Cold drawn

430

62

360

52.5

15

40

121

1025

G10250

Hot rolled

400

58

220

32

25

50

116

Cold drawn

440

64

370

54

15

40

126

1026 1030 1035 1037 1038

G10260 G10300 G10350 G10370 G10380

Hot rolled

440

64

240

35

24

49

126

Cold drawn

490

71

410

60

15

40

143

Hot rolled

470

68

260

37.5

20

42

137

Cold drawn

520

76

440

64

12

35

149

Hot rolled

500

72

270

39.5

18

40

143

Cold drawn

550

80

460

67

12

35

163

Hot rolled

510

74

280

40.5

18

40

143

Cold drawn

570

82

480

69

12

35

167

Hot rolled

520

75

280

41

18

40

149

Cold drawn

570

83

480

70

12

35

163

1039

G10390

Hot rolled

540

79

300

43.5

16

40

156

Cold drawn

610

88

510

74

12

35

179

1040

G10400

Hot rolled

520

76

290

42

18

40

149

Cold drawn

590

85

490

71

12

35

170

1042

G10420

Hot rolled

550

80

300

44

16

40

163

1043

G10430

65 70 70 65 70 70 65 65 75 70 65 65 65 60 60

Cold drawn

610

89

520

75

12

35

179

60

NCD

590

85

500

73

12

45

179

70

Hot rolled

570

82

310

45

16

40

163

Cold drawn

630

91

530

77

12

35

179

60 70

NCD

600

87

520

75

12

45

179

1044

G10440

Hot rolled

550

80

300

44

16

40

163

1045

G10450

Hot rolled

570

82

310

45

16

40

163

Cold drawn

630

91

530

77

12

35

179

55

ACD

590

85

500

73

12

45

170

65

1046

1049

1050

1055

G10460

G10490

G10500

G10550

Hot rolled

590

85

320

47

15

40

170

Cold drawn

650

94

540

79

12

35

187

55

ACD

620

90

520

75

12

45

179

65

Hot rolled

600

87

330

48

15

35

179

Cold drawn

670

97

560

81.5

10

30

197

45

ACD

630

92

530

77

10

40

187

55

Hot rolled

620

90

340

49.5

15

35

179

Cold drawn

690

100

580

84

10

30

197

45

ACD

660

95

550

80

10

40

189

55

Hot rolled

650

94

360

51.5

12

30

192


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01 Sep 2005

ACD

660

96

560

81

10

40

197

54

12

30

201

1060

G10600

Hot rolled

680

98

370

SACD

620

90

480

70

10

45

183

1064

G10640

Hot rolled

670

97

370

53.5

12

30

201

SACD

610

89

480

69

10

45

183

1065

G10650

Hot rolled

690

100

380

55

12

30

207

SACD

630

92

490

71

10

45

187

1070

G10700

Hot rolled

700

102

390

56

12

30

212

SACD

640

93

500

72

10

45

192

1074

G10740

Hot rolled

720

105

400

58

12

30

217

SACD

650

94.5

500

73

10

40

192

1078

G10780

Hot rolled

690

100

380

55

12

30

207

SACD

650

94

500

72.5

10

40

192

1080

G10800

Hot rolled

770

112

420

61.5

10

25

229

SACD

680

98

520

75

10

40

192

1084 1085 1086 1090 1095

G10840 G10850 G10860 G10900 G10950

Hot rolled

820

119

450

65.5

10

25

241

SACD

690

100

530

77

10

40

192

Hot rolled

830

121

460

66.5

10

25

248

SACD

690

100.5

540

78

10

40

192

Hot rolled

770

112

420

61.5

10

25

229

SACD

670

97

510

74

10

40

192

Hot rolled

840

122

460

67

10

25

248

SACD

700

101

540

78

10

40

197

Hot rolled

830

120

460

66

10

25

248

SACD

680

99

520

76

10

40

197

55 60 60 60 55 55 55 45 45 45 45 45 45

Manganese, maximum >1.00% 1524 1527 1536 1541

1548

1552

G15240 G15270 G15360 G15410

G15480

G15520

Hot rolled

510

74

280

41

20

42

149

Cold drawn

570

82

480

69

12

35

163

Hot rolled

520

75

280

41

18

40

149

Cold drawn

570

83

480

70

12

35

163

Hot rolled

570

83

310

45.5

16

40

163

Cold drawn

630

92

530

77.5

12

35

187

60 65 55

Hot rolled

630

92

350

51

15

40

187

Cold drawn

710

102.5

600

87

10

30

207

45

ACD

650

94

550

80

10

45

184

60

Hot rolled

660

96

370

53

14

33

197

Cold drawn

730

106.5

620

89.5

10

28

217

45

ACD

640

93.5

540

78.5

10

35

192

50

Hot rolled

740

108

410

59.5

12

30

217

ACD 680 98 570 83 10 40 193 50 (a) ACD, annealed cold drawn; NCD, normalized cold drawn; SACD, spheroidized annealed cold drawn. (b) Cold drawn 1212 = 100%. Source: Ref 8

Table 10 Estimated mechanical properties and machinability ratings of resulfurized carbon steel bars All 1100 and 1200 series steels are rated on the basis of 0.10% maximum or coarse-grain melting practice. Estimated minimum values Steel designation SAE and/or AISI No.

UNS No.

Tensile strength

Type of processing

MPa


ksi

Yield strength

MPa

ksi

Elongation Reduction in 50 mm in (2 in.), % area, %

Average machinabi Hardness, lity HB rating(a)

Page 439


1108

G11080

1117

G11170

1132

G11320

1137

G11370

1140

G11400

1141

G11410

1144

G11440


01 Sep 2005

Hot rolled

340

50

190

27.5

30

50

101

Cold drawn

390

56

320

47

20

40

121

Hot rolled

430

62

230

34

23

47

121

Cold drawn

480

69

400

58

15

40

137

Hot rolled

570

83

310

45.5

16

40

167

Cold drawn

630

92

530

77

12

35

183

Hot rolled

610

88

330

48

15

35

179

Cold drawn

680

98

570

82

10

30

197

Hot rolled

540

79

300

43.5

16

40

156

Cold drawn

610

88

510

74

12

35

170

Hot rolled

650

94

360

51.5

15

35

187

Cold drawn

720

105.1

610

88

10

30

212

Hot rolled

670

97

370

53

15

35

197

Cold drawn

740

108

620

90

10

30

217

47

15

40

170

1146

G11460

Hot rolled

590

85

320

Cold drawn

650

94

550

80

12

35

187

1151

G11510

Hot rolled

630

92

340

50.5

15

35

187

Cold drawn

700

102

590

86

10

30

207

1211

G12110

Hot rolled

380

55

230

33

25

45

121

Cold drawn

520

75

400

58

10

35

163

1212

G12120

Hot rolled

390

56

230

33.5

25

45

121

Cold drawn

540

78

410

60

10

35

167

1213

G12130

Hot rolled

390

56

230

33.5

25

45

121

Cold drawn

540

78

410

60

10

35

167

12L14

G12144

Hot rolled

390

57

230

34

22

45

121

540

78

410

60

10

35

163

Cold drawn (a) Cold drawn 1212 = 100%. Source: Ref 8

80 90 75 70 70 70 80 70 65 95 100 135 170

Table 11 Hardness and machinability ratings of cold-drawn alloy steels Steel designation AISI and/or SAE No.

UNS No.

Machinabil ity rating(a)

1330

G13300

55

1335

G13350

1340

G13400

1345 4023

Range of typical hardness, HB

Microstructur e type(b)

Annealed and cold drawn

179/235

A

55


179/235

A

50


183/241

A

G13450

45


183/241

A

G40230

70

Cold drawn

156/207

C

4024

G40240

75

Cold drawn

156/207

C

4027

G40270

70


167/212

A

4028

G40280

75


167/212

A

4032

G40320

70


174/217

A

4037

G40370

70


174/217

A

4042

G40420

65


179/229

A

4047

G40470

65


179/229

A

4118

G41180

60

Cold drawn

170/207

C

4130

G41300

70


187/229

A

4135

G41350

70


187/229

A

4137

G41370

70


187/229

A

4140

G41400

65


187/229

A

Condition


Page 440



01 Sep 2005

4142

G41420

65


187/229

A

4145

G41450

60


187/229

A

4147

G41470

60


187/235

A

4150

G41500

55


187/241

A, B

4161

G41610

50

Spheroidized and cold drawn

187/241

B, A

4320

G43200

60


187/229

D, B, A

4340

G43400

50


187/241

B, A

E4340

G43406

50


187/241

B, A

4422

G44220

65

Cold drawn

170/212

A

4427

G44270

65


170/212

A

4615

G46150

65

Cold drawn

174/223

C

4617

G46170

65

Cold drawn

174/223

C

4620

G46200

65

Cold drawn

183/229

C

4626

G46260

70

Cold drawn

170/212

C

4718

G47180

60

Cold drawn

187/229

C

4720

G47200

65

Cold drawn

187/229

C

4815

G48150

50


187/229

D, B

4817

G48170

50


187/229

D, B

4820

G48200

50


187/229

D, B

50B40

G50401

65


174/223

A

50B44

G50441

65


174/223

A

5046

G50460

60


174/223

A

50B46

G50461

60


174/223

A

50B50

G50501

55


183/235

A

5060

G50600

55

Spheroidized annealed and cold drawn

170/212

B

50B60

G50601

55


170/212

B

5115

G51150

65

Cold drawn

163/201

C

5120

G51200

70

Cold drawn

163/201

C

5130

G51300

70


174/212

A

5132

G51320

70


174/212

A

5135

G51350

70


179/217

A

5140

G51400

65


179/217

A

5147

G51470

65


179/229

A

5150

G51500

60


183/235

A, B

5155

G51550

55


183/235

A, B

5160

G51600

55


179/217

B

51B60

G51601

55


179/217

B

50100

G50986

40


183/241

B

51100

G51986

40


183/241

B

52100

G52986

40


183/241

B

6118

G61180

60

Cold drawn

179/217

C

6150

G61500

55


183/241

B, A

8115

G81150

65

Cold drawn

163/202

C

81B45

G81451

65


179/223

A

8615

G86150

70

Cold drawn

179/235

C

8617

G86170

70

Cold drawn

179/235

C

8620

G86200

65

Cold drawn

179/235

C

8622

G86220

65

Cold drawn

179/235

C

8625

G86250

60


179/223

A


Page 441



01 Sep 2005

8627

G86270

60


179/223

A

8630

G86300

70


179/229

A

8637

G86370

65


179/229

A

8640

G86400

65


184/229

A

8642

G86420

65


184/229

A

8645

G86450

65


184/235

A

86B45

G86451

65


184/235

A

8650

G86500

60


187/248

A, B

8655

G86550

55


187/248

A, B

8660

G86600

55


179/217

B

8720

G87200

65

Cold drawn

179/235

C

8740

G87400

65


184/235

A

8822

G88220

55

Cold drawn

179/223

B

9254

G92540

45


187/241

B

9260

G92600

40


184/235

B

9310

G93106

50


184/229

D

94B15

G94151

70

Cold drawn

163/202

C

94B17

G94171

70

Cold drawn

163/202

C

94B30 G94301 70 Annealed and cold drawn 170/223 A (a) Based on cutting with high-speed tool steels and a machinability rating of 100% for 1212 steel. (b) Type A is predominantly lamellar pearlite and ferrite. Type B is predominantly spheroidized. Type C is a hot-rolled structure that depends on grade, size, and rolling conditions of the producing mill. The structure may be coarse or fine pearlite or bainite. The pearlite at low magnification may be blocky or acicular. For descriptive information, see Ref 9. Type D is a structure resulting from a subcritical anneal or temper anneal. It is usually a granular or spheroidized carbide condition confined to the hot-rolled grain pattern, which may be blocky or acicular. Source: Ref 8

Strength Considerations Cold drawing increases the tensile and yield strengths of hot-rolled carbon steel bars by about 10 and 70%, respectively. For example, hot-rolled low-carbon steel bars with yield-to-tensile-strength ratios of about 0.55 will have ratios of about 0.85 after cold drawing. Elongation, reduction in area, and impact strength will be decreased, but these changes are relatively insignificant in many structural and engineering applications. Improvements in strength resulting from cold drawing are of interest to the design engineer seeking a better strength-to-weight ratio or a reduction in costs by elimination of alloy additions and heat treatment. They may also be useful in applications involving threads, notches, cutouts, size stability, and other design factors that might affect strength adversely. It should be remembered, however, that the spread in mechanical properties within a single grade of cold-drawn steel may not lend itself to as close control as may be obtained by heat treating to the same strength level. Some variation in properties may result from differences in mill technique and amount of cold work. It should also be remembered that turned and polished bars and turned, ground, and polished bars have the same mechanical properties as hot-rolled bars. Tensile property values that may be expected for cold-drawn carbon steel bars processed with normal drafts are shown in Fig. 14 , which covers bars from 19 to 32 mm (3=4to 11=4in.) in diameter. For smaller bars, tensile and yield strengths will be slightly higher; for larger bars, they will be slightly lower. The values of yield strength shown in Fig. 14 are conservative and are very close to the expected minimums. Fig. 14 Effect of cold drawing on the tensile properties of plain carbon and resulfurized carbon steels. Comparison of tensile strength and yield strength of 25 mm (1 in.) diam cold-drawn (a) plain carbon steels and (b) resulfurized carbon steels. Comparison of yield strengths of 25 mm (1 in.) diam hot-rolled and 25 mm (1 in.) diam cold-drawn (c) plain carbon and (d) resulfurized carbon steels. Source: Ref 10


Page 442



01 Sep 2005

Figures 15 , 16 , 17 , 18 , and 19 present actual tensile test data for five grades of steel that are commonly cold drawn. In these charts, most of the minimum values for tensile and yield strengths are considerably higher than those shown in Fig. 14 . Numerous factors may significantly affect these values, the principal ones being: • The variation in composition from the low to the high limit of the specification. For example, in 1045 steel, carbon may vary from 0.40 to 0.53%, and manganese from 0.57 to 0.93%, on the basis of bar-to-bar checks (heat range plus allowable variations in product analysis) • The effects of solidification rate, type of solidification pattern, and segregation • Temperature of the bar in the last pass in the the hot mill and rate of cooling after rolling. These factors affect hardness, particularly when carbon content is 0.30% or more. With the same draft, bars of higher hardness will show higher strength and lower ductility after cold drawing than will bars of lower hardness • Hot roll reduction • Macro- and microstructure. Finish temperature has a dramatic effect on grain size


Page 443



01 Sep 2005

Therefore, the minimum property values for conventionally cold-drawn material must be conservatively low. Fig. 15 Distribution of tensile properties of grade 1018 cold-drawn steel bars. Tested were round bars, 19 to 32 mm (3=4to 11=4in.) in diameter, from 58 heats from 5 mills.

Fig. 16 Distribution of tensile properties of grade 1045 cold-drawn steel bars. Tested were round bars, 19 to 32 mm (3=4to 11=4in.) in diameter, from 40 heats from 5 mills.

Fig. 17 Distribution of tensile properties of grade 1117 cold-drawn steel bars. Tested were round bars, 25 mm (1 in.) in diameter, from 25 heats from 2 mills.


Page 444



01 Sep 2005

Fig. 18 Distribution of tensile properties of grade 1137 cold-drawn steel bars. Tested were round bars, 25 mm (1 in.) in diameter, from 25 heats from 2 mills.

Fig. 19 Distribution of tensile properties of grade 12L14 cold-drawn steel bars. Tested were round bars, 19 to 38 mm (3=4to 11=2in.) in diameter, and hexagon bars, 14 to 25 mm (9=16to 1 in.) in diameter, from 64 heats from 1 mill.


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01 Sep 2005

Special Die Drawing Two special die-drawing processes have been developed to give improved properties over those offered by standard drawing practices. These processes are cold drawing using heavier-than-normal drafts, followed by stress relieving; and drawing at elevated temperatures. In the production of steel bars by these special processes, drafts of approximately 10 to 35% reduction in cross-sectional area are employed at room or elevated temperature, depending on the practices and facilities of the individual producer. Stress-relieving temperatures vary over a similarly wide range, depending on producer facilities and end-product requirements. Typical tensile properties of plain carbon and alloy steel bars of medium carbon content subjected to special die-drawing processing are given in Table 12 . Table 12 Typical mechanical properties of special-die-drawn carbon and alloy steel bars of medium-carbon content Tensile strength

Bar size range

Steel grades

mm

Yield strength

Elong ation in 50 Reductio n mm in Hardness (2 area, , in.), %(a) HB(b) %(a)

in.

MPa

ksi

MPa

ksi

Up to 76 round

Up to 3 round

825

120

690

100

10.0

25.0

241−321

1052, 1141

6.4 through 89 round

1

825

120

690

100

10.0

25.0

241−321

1144, 1151

6.4 through 114 round and 6.4 through 51 hexagon

1

=4 through 41=2 round and 1=4 through 2 hexagon

825

120

690

100

10.0

25.0

248−321

1144

6.4 through 6.4 round and 6.4 through 38 hexagon

1

=4 through 21=2 round and 1=4 through 11=2 hexagon

965

140

860

125

5.0

15.0

280

41xx, 51xx(c)

11.1 through 89 hexagon

7

860

125

725

105

14.0

45.0

269

41xx(c), 51xx(c)


7

=16 through 31=2 hexagon

103.5

150

895

130

10.0

35.0

302

41xx(c)


7

1170

170

1140

155

5.0

20.0

355

Carbon steels 1541, 1045

1

=4 through 3 =2 round

Alloy steels =16 through 31=2 hexagon

=16 through 31=2 hexagon


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01 Sep 2005

(a) Typical minimum. (b) Typical hardness range, subject to negotiation. Hardness is taken on a flat below decarburization or on the midradius. In case of a disagreement between hardness and tensile or yield strength, the latter properties govern. (c) May contain lead or tellurium or other additives for improved machinability. Source: Ref 8

Heavy Drafts. Because of the engineering and economic advantages of cold-finished steels, a considerable effort has been made to improve the uniformity of mechanical properties after cold drawing. This has been accomplished by using heavier-than-normal drafts (10 to 35% reduction) followed by subcritical annealing. Stepwise or tandem drawing has been resorted to in order to avoid the formation of internal transverse fissures (cupping or bambooing) that may result from heavy drafts taken in one pass. The trend of property improvement resulting from this special practice can be seen in the comparison of normal and heavy drafts shown in Fig. 20 . Fig. 20 Effect of draft and stress-relieving temperature on the tensile properties of cold-drawn carbon steel bars. Solid curves are for bars given a normal draft; dashed curves are for bars given a heavy draft.

Heavier drafts produce higher tensile and yield strengths. Elongation can be substantially improved by stress relieving at 510 °C (950 °F), and this treatment still provides tensile and yield strengths higher than those obtainable with normal drafts. The combination of properties resulting from heavier drafts and higher stress-relieving temperatures is most desirable from the design standpoint. Such processing is most effective when applied to medium-carbon steels of either normal or higher manganese content. In the medium-manganese range, 1045 and 1050 respond most favorably; 1137, 1141, 1144, 1527, 1536, and 1541 show good response for the higher-manganese grades. The sulfur content of the 11xx steels improves machinability without lowering mechanical properties. Typical tensile properties for 1144 bars subjected to normal and heavy drafts are given in Table 13 . Table 13 Typical tensile properties of cold-drawn and stress-relieved 1144 grade (UNS G11440) carbon steel bars subjected to normal and heavy drafts Size Diameter of round, thickness of square, or distance between parallel faces of hexagon or flat mm

in.

Strength

Round or hexagon(a) mm

in.

Tensile strength minimum MPa

ksi

Yield strength minimum MPa

ksi

Elongation in 50 mm (2 in.) minimum, %

Reduction in area minimum, %

Normal draft To 22 incl

To 7=8 incl

...

...

725

105

655

95

10

30

>22−32 incl

> =8−1 =4 incl

...

...

690

100

620

90

10

30

>32−51 incl

>11=4−2 incl

...

...

655

95

585

85

10

25

>51−76 incl

>2−3 incl

...

...

620

90

550

80

10

20

...

...

585

85

520

75

10

20

>76−114 incl

7

1

1

>3−4 =2 incl

Heavy draft ... ... ...

... ... ...

To 22 incl

To 7=8 incl

795

115

690

100

8

25

>22−32 incl

7

1

=8−1 =4 incl

795

115

690

100

8

25

>32−51 incl

1 =4−2 incl

795

115

690

100

8

25


1

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...

...


>51−76 incl

01 Sep 2005

2−3 incl

795

115

690

100

8

20

1

795

115

690

100

7

20

... ... >76−114 incl 3−4 =2 incl incl, inclusive.(a) Maximum size for hexagons is 38 mm (11=2in.)

Steel bars that have been cold drawn using heavier-than-normal drafts and then stress relieved are often used in place of quenched and tempered bars, and as already noted, several resulfurized grades (1137, 1141, and 1144) respond readily to this process with resulting high strength values. Because the microstructures of these steels are still pearlitic, they machine more easily than their quenched and tempered counterparts. Therefore, although these grades cost more than nonresulfurized grades, they can provide significant savings in manufacturing costs, chiefly through the elimination of heat treating. However, even through the strength of the special cold-drawn and stress-relieved bars may be equal to that of quenched and tempered steel, it is not advisable to translate other properties from one condition to the other. The torsional strength and endurance limits of these special-process bars are similar to those of quenched and tempered bars at the same strength level. The same is true for the wear resistance of bars of the same surface hardness. The impact test values of the process bars, as measured by Izod or Charpy notched-bar test, are lower than those of quenched and tempered carbon steel bars and are significantly lower than those of quenched and tempered alloy steel bars. Failures of machine components usually result from fatigue, corrosion, wear, or shock loading. With the possible exception of the latter, there is no known correlation between instances of failure and the notched-bar impact test. When low temperatures or high pressures are involved and where there is doubt as to the suitability of these special-process bars, the design of the part should be reviewed. Drawing at elevated temperatures between 95 and 540 °C (200 and 1000 °F), a special proprietary process, can, under optimum conditions, produce steel bars that have higher tensile and yield strengths than those of bars cold drawn with the same degree of reduction. The relative effects of cold drawing followed by stress relieving and of drawing at elevated temperature can be seen in Fig. 21 . Both processes were carried out with 20% draft on 25 mm (1 in.) diam bars of 1144 steel. As shown in Fig. 21 , elevated-temperature drawing affects tensile strength considerably, giving values somewhat greater than those for cold-drawn stress-relieved bars. For yield strength, the same general effects are evident, but the difference between the two processes is not as pronounced. The elongation values are quite similar for both processes. Fig. 21 Effects of stress-relieving or drawing temperature on the (a) tensile strength and (b) yield strength of cold-drawn and stress-relieved bars and on hot drawn bars of 1144 steel. Bars, all from the same heat of steel and approximately 25 mm (1 in.) in diameter before drawing, were subjected to a draft of about 20%. Source: Ref 11

Figures 22 , 23 , 24 , and 25 show the effects of two drafts and increasing drawing temperatures on bars of each of four steels. Between 260 and 315 °C (500 and 600 °F), tensile and yield strengths reach maximums and elongation and reduction in area reach minimums. Strength for any given drawing temperature increases with increased draft, with minor exceptions. Depending on chemical composition, a certain minimum draft is required to fully develop the effect on strength properties of drawing at elevated temperature. In general, the draft taken at room temperature would have to be doubled in order to match the strength developed in drawing at elevated temperature. The yield strengths of all four steels were increased from 12 to 35%, depending on percentage reduction in drawing and on drawing temperature. Carbon steels were less affected by increased draft than alloy steel 4140 when drawn at elevated temperatures. Fig. 22 Effects of drawing temperature and percentage reduction on mechanical properties of 19 mm (3=4in.) diam cold-drawn 1018 steel bars. Source: Ref 12


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Fig. 23 Effects of drawing temperature and percentage reduction on mechanical properties of 14 mm (9=16in.) diam cold-drawn 1080 steel bars. Source: Ref 12



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01 Sep 2005


The pronounced effect of drawing at elevated temperatures changes the shape of the stress-strain curve from round to sharp-kneed, as shown in Fig. 26 . Hot drawing reverses the effect of cold work, that is, automatically stress relieves the steel. This effect on the stress-strain curve is significant in applications in which minimum plastic deformation is permissible, such as a stud that requires a proof load high in relation to its tensile strength. Fig. 26 Effect of drawing temperature on the shape of the tensile stress-strain curve for hot-drawn 1144 steel bars. Bars were reduced 7.2% to a diameter of 21.4 mm (27=32in.). Source: Ref 12


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Typical yield strength minimums for 1144 bars before and after drawing are as follows: MPa

ksi

As-hot-rolled

345

50

Cold drawn (normal draft) and stress relieved at low temperature

550

80

Cold drawn (heavy draft) and stress relieved at high temperature

690

100

Drawn at elevated temperature

860

125

The effect of drawing temperature on notched-bar impact values is shown in Fig. 27 . Although impact values decrease with the first increase in drawing temperature, they later rise significantly. Thus, it may be possible to select a drawing temperature that will provide both good static strength and satisfactory notched-bar impact strength. Fig. 27 Effect of drawing temperature on the impact properties of hot-drawn bars of two steels. Source: Ref 12

REFERENCES 1. J.G. Bralla, Handbook of Product Design for Manufacturing, McGraw-Hill, 1986 2. Alloy, Carbon, and High Strength Low Alloy Steels, Semifinished for Forging; Hot Rolled Bars; Cold Finished Steel Bars; Hot Rolled Deformed and Plain Concrete Reinforcing Bars, AISI Steel Products Manual, American Iron and Steel Institute, 1986 3. Handbook of Machining Data for Cold Finished Steel Bars, LTV Steel Flat Rolled and Bar Company, 1985 4. Steel⎯Bars, Forgings, Bearing, Chain, Springs, Vol 1.05, Annual Book of ASTM Standards, American Society for Testing and Materials, 1989 5. L.J. Ebert, Report WAL 310/90-85 to Watertown Arsenal, 1955 6. H. Buhler and H. Bucholz, Influence of Cold Drawn Reduction Upon Stresses in Round Bars, Arch. Eisenhüttenwes., Vol 7, 1934, p 427−430 7. E. Dieter, Mechanical Metallurgy, McGraw-Hill, 1976


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8. Materials, Vol 1, 1989 SAE Handbook, Society of Automotive Engineers, 1989 9. "U.S. Air Force Machinability Report," Vol 2, Curtiss-Wright Corporation, 1951 10. "Estimated Properties and Machinability of Plain Carbon and Re-sulfurized Plain Carbon Steel Bars," SAE J414, SAE Handbook Supplement HS30, Society of Automotive Engineers, 1976, p 3.14 11. E.S. Nachtman and E.B. Moore, J. Met., April 1955 12. E.S. Nachtman and E.B. Moore, J. Met., April 1958


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Steel Wire Rod

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Carbon and Low-Alloy Steels Steel Wire Rod Revised by R.J. Glodowski, Armco, Inc. WIRE ROD is a semifinished product rolled from billet on a rod mill and is used primarily for the manufacture of wire. The steel for wire rod is produced by all the modern processes, including the basic oxygen, basic open hearth, and electric furnace processes. Steel wire rod is usually cold drawn into wire suitable for further drawing; for cold rolling, cold heading, cold upsetting, cold extrusion, or cold forging; or for hot forging. Although wire rod may be produced in several regular shapes, most is round in cross section. Round rod is usually produced in nominal diameters of 5.6 to 18.7 mm (7=32to 47=64in.), advancing in increments of 0.4 mm (1=64in.).* [*Because steel wire rod manufactured in the United States is customarily produced to fractional-inch sizes, rather than decimal-inch or millimeter sizes, the millimeter conversion for wire rod sizes may not be a multiple of the 0.4 mm (1=64in.) increment size.] As the rod comes off the rolling mill, it is formed into coils. These coils are usually about 760 mm (30 in.) in inside diameter and weigh up to 2000 Kg (4400 lb). The dimensions and maximum weight of a single coil are determined by the capabilities of the rolling mill. Coil weights that exceed the capabilities of the rolling mill sometimes can be obtained by welding two or more coils together. The standard tolerances are ±0.4 mm (±1=64in.) on the diameter and 0.64 mm (0.025 in.) maximum out-of-roundness. Producers of wire rod may market their product as rolled, as cleaned and coated, or as heat treated, although users generally prefer to do such preparations themselves. These operations are explained in the following sections, along with the several recognized quality and commodity classifications applicable to steel wire rods.

Cleaning and Coating Mill scale is cleaned from steel wire rods by pickling or caustic cleaning followed by water rinsing, or by mechanical means such as shot blasting with abrasive particles or reverse bending over sheaves. The chemical cleaning of steel wire rods is always followed by a supplementary coating operation. Lime, borax, or phosphate coating is applied to provide a carrier for the lubricant necessary for subsequent processing into wire. In lime coating, practices may be varied in order to apply differing amounts of lime on the rods depending on the customer requirements. Phosphate-coated rods may have a supplementary coating of lime, borax, or water-soluble soap. Mechanically descaled rods may be drawn without coating using only wire drawing soaps, or may be coated in a fashion similar to that used for chemically cleaned rods.

Heat Treatment The heat treatments commonly applied to steel wire rod, either before or during processing into wire, include annealing, spheroidize annealing, patenting, and controlled cooling. Annealing commonly involves heating to a temperature near or below the lower critical temperature and holding at that temperature for a sufficient period of time, followed by slow cooling. This process softens the steel for further processing, but not to the same degree as does spheroidize annealing. Spheroidize annealing involves prolonged heating at a temperature near or slightly below the lower critical temperature (or thermal cycling at about the lower critical temperature), followed by slow cooling, with the object of changing the shape of carbides in the microstructure to globular (spheroidal), which produces maximum softness. Patenting is a heat treatment usually confined to medium-high-carbon and high-carbon steels. In this process, individual strands of rod or wire are heated well above the upper critical temperature and then are cooled comparatively rapidly in air, molten salt, molten lead, or a fluidized bed. The object of patenting is to develop a microstructure of homogeneous, fine pearlite. This treatment generally is employed to prepare the material for subsequent wire drawing. Controlled cooling is a heat treatment performed in modern rod mills in which the rate of cooling after hot rolling is carefully controlled. The process imparts uniformity of properties and some degree of control over scale, grain size, and microstructure.

Carbon Steel Rod Carbon steels are those steels for which no minimum content is specified or required for chromium, nickel, molybdenum, tungsten, vanadium, cobalt, niobium, titanium, zirconium, aluminum, or any other element added to obtain a desired alloying effect; for which specified minimum copper content does not exceed 0.40%; for which specified maximum manganese content does not exceed 1.65%; and for which specified maximum silicon or copper content does not exceed 0.60%. In all carbon steels,


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small quantities of certain residual elements, such as chromium, nickel, molybdenum, and copper, are unavoidably retained from raw materials. These elements are considered incidental, although maximum limits are commonly specified for specific end-uses. Carbon steel rods are produced in various grades, or compositions: • • • •

Low-carbon steel wire rods (maximum carbon content ≤0.15%) Medium-low-carbon steel wire rods (maximum carbon content >0.15%, but ≤0.23%) Medium-high-carbon steel wire rods (maximum carbon content >0.23%, but ≤0.44%) High-carbon steel wire rods (maximum carbon content >0.44%)

Ordinarily, sulfur and phosphorus contents are kept within the usual limits for each grade of steel, while carbon, manganese, and silicon contents are varied according to the mechanical properties desired. Occasionally, sulfur and/or phosphorus may be added to the steel to improve the machinability.

Qualities and Commodities of Carbon Steel Rod Rod for the manufacture of carbon steel wire is produced with manufacturing controls and inspection procedures intended to ensure the degree of soundness and freedom from injurious surface imperfections necessary for specific applications. The various quality descriptors and commodities applicable to carbon steel wire rod are described below. Industrial quality rod is manufactured from low-carbon or medium-low-carbon steel and is intended primarily for drawing into industrial quality wire. Rod of this quality is available in the as-rolled or heat-treated conditions. Practical limitations for drawing are: low-carbon rod 5.6 mm (7=32in.) in diameter can be drawn without intermediate annealing to 2.0 mm (0.080 in.) by five conventional drafts; medium-low-carbon rod 5.6 mm (7=32in.) in diameter can be drawn without intermediate annealing to 2.69 mm (0.106 in.) by four conventional drafts. Chain Quality Rod. Rod for the manufacture of wire to be used for resistance welded chain is made from low-carbon and medium-low-carbon steel produced by practices that ensure their suitability for drawing into wire for this end-use. Good butt welding uniformity characteristics and internal soundness are essential for this application. Rod for the manufacture of wire to be used for fusion welded chain can be produced from specially selected low-carbon rimmed steel, but is more often made from continuous cast steel. Fine wire quality rod is suitable for drawing into small-diameter wire either without intermediate annealing treatments or with only one such treatment. Rod 5.6 mm (7=32in.) in diameter can be direct drawn into wire as fine as 0.9 mm (0.035 in.) without intermediate annealing. Wire finer than 0.9 mm (0.035 in.), for such products as insect-screen wire, weaving wire, and florist wire, is usually drawn in two steps: reducing to an intermediate size no smaller than 0.9 mm (0.035 in.), followed by annealing and redrawing to final size. Fine wire quality rod is generally rolled from steel of grade 1005 or 1006 produced using techniques to provide good surface finish and internal cleanliness. In addition to these precautions, the producer may subject the rod to tests such as fracture or macroetch tests. Cold finishing quality rod is intended for drawing into cold finished bars; the manufacture of such rod is controlled to ensure suitable surface conditions. Heading, Cold Extrusion, or Cold Rolling Quality Rod. Rod used for the manufacture of heading, forging, cold extrusion or cold rolling quality wire is produced by closely controlled manufacturing practices. It is subject to mill testing and inspection to ensure internal soundness and freedom from injurious surface imperfections. Heat treatment as a part of wire mill processing is very important in the higher carbon grades of steel. For common upsetting, represented by the production of standard trimmed hexagon-head cap screws, 1016 to 1038 steel wire drawn from annealed rod is suitable. Wire for moderate upsetting, also produced from 1016 to 1038 steel, should be drawn from spheroidize annealed rod or should be in-process annealed. Wire for severe heading and forging, produced from rod of 1016 to 1541 steel, should be spheroidize annealed in process or at finished size. Rod of this quality is not intended for recessed-head or similar special-head applications. In the production of rod for heading, forging, or cold extrusion is killed carbon steels with nominal carbon contents of 0.16% or more (AISI grades 1016 or higher), both austenitic grain size and decarburization should be controlled. Such steels can be produced with either fine or coarse austenitic grains, depending on the type of heat treatment and end-use. The maximum allowable amounts of decarburization as defined by the average value for the depth of the layer of free ferrite plus the layer of partial decarburization (the total affected depth) and the average depth of the layer of free ferrite alone are given below: Average depth of decarburization Free ferrite

Total affected

Nominal rod diameter, mm (in.)

mm

in.

mm

in.

≤9.6 (≤ =8)

0.10

0.004

0.30

0.012

3

10−12.8 ( =64− =2)

0.13

0.005

0.36

0.014

13−18.7 ( =64− =64)

0.15

0.006

0.41

0.016

25 33

1

47

If decarburization limits closer than these standard limits are required in a given manufactured product, it is sometimes


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Steel Wire Rod

01 Sep 2005

necessary to incorporate means for carbon restoration in the annealing process. When there are discrepancies in decarburization test results, it is customary to make heat-treatment tests of the finished product to determine suitability for the particular application. Wood screw quality rod includes low-carbon resulfurized and nonresulfurized wire rod for drawing into wire for the manufacture of slotted-head screws only, not for recessed-head or other special-head screws. Scrapless Nut Quality Rod. Rod to be drawn into wire for scrapless nuts is produced by specially controlled manufacturing practices. It is subjected to mill tests and inspection designed to ensure internal soundness; freedom from injurious segregation and injurious surface imperfections; and satisfactory performance during cold heading, cold expanding, cold punching, and thread tapping. Rod for scrapless nut wire commonly is made from low-carbon, resulfurized steels. Nonresulfurized steels are also used; these steels ordinarily are furnished only in grades containing more carbon than the resulfurized grades and with phosphorus content not exceeding 0.035% and sulfur content not exceeding 0.045% by heat analysis. In making resulfurized steel for scrapless nut quality rod, either an ingot or continuous casting process can be used. In an ingot manufacturing process, sometimes the sulfur content is obtained through delayed mold additions to a conventional nonresulfurized rimming steel. The purpose of such a practice is to produce a steel consisting of a rim of low-sulfur steel suitable for expansion during nut forming around a high-sulfur interior section suitable for the piercing and threading operations involved in making scrapless nuts. When high sulfur content is secured through such mold additions, sulfur analyses are made on the solid billets rather than by heat analysis. It is customary to produce these steels to a specified sulfur range of either 0.08 to 0.13% or 0.04 to 0.09%. Because of the practice used in making the steel and the degree to which sulfur segregates, the sulfur content at various locations in a billet may vary from the indicated range. When a continuous casting process is used to make resulfurized steel, the sulfur content is typically more uniform than the ingot process. However, continuous casting precludes the rimming practice described in the above paragraph. Severe cold heading, severe cold extrusion, or severe scrapless nut quality rod is used for severe single-step or multiple-step cold forming where intermediate heat treatment and inspection are not possible. Rod of this quality is produced with carefully controlled manufacturing practices and rigid inspection practices to ensure the required degree of internal soundness and freedom from surface imperfections. A fully killed fine-grain steel is usually required for the most difficult operations. Normally, the wire made from this quality rod is spheroidize annealed, either in process or after drawing finished sizes. Decarburization limits and the steels to which they apply are the same as those described in the section "Heading, Cold Extrusion, or Cold Rolling Quality Rod" in this article. Welding-quality rod is used to make wire for gas or electric-arc welding filler metal. Welding-quality rod can be made from selected ingots or billets of low-carbon rimmed, capped, or killed steel, but is preferably made from continuous cast steel. It is produced to several restricted ranges and limits to chemical composition; an example of the restricted ranges and limits for low-carbon, arc welding wire rod is shown below: Element

Heat analysis, %

Carbon

0.10−0.15

Manganese

0.40−0.60

Phosphorus

0.025 max

Sulfur

0.035 max

Silicon

0.030 max

Rod for welding-quality wire constitutes an exception to the general practice that rimmed or capped steel is not commonly subject to product analysis. Experience to date has shown the necessity for close control of composition, and therefore only billets from those portions of the ingot that conform to the applicable ranges and limits are used for welding-quality rod. For the majority of welding-quality rod that is made from continuous cast steel, these product checks may not be necessary. Medium-high-carbon and high-carbon quality rod is wire rod intended for drawing into such products as strand wire, lockwasher wire, tire bead wire, upholstery spring wire, rope wire, screen wire (for heavy aggregate screens), aluminum cable steel reinforced core wire, and prestressed concrete wire. These wire qualities are normally drawn directly from patented or control-cooled rod. When drawing to sizes finer than 2.0 mm (0.080 in.) (from 5.6 mm, or 7=32in., rod), it is customary to employ in-process heat treatment before drawing to finish size. Medium-high-carbon and high-carbon quality rod is not intended for the manufacture of higher-quality wires such as music wire or valve spring wire. Rod for Special Purposes. In addition to the carbon steel rod commodities described above, which have specific quality descriptors, several other commodities are produced, each having the characteristics necessary for a specific application, but for which no specific quality descriptor exists. Some of these commodities are made to standard specifications; the others are made to proprietary specifications that are mutually acceptable to both producer and user. Rod for music wire, valve spring wire, and tire cord wire is rolled and conditioned to ensure the lowest possible incidence of imperfections. Surface imperfections are objectionable because they lower the fatigue resistance that is important in many of the end products made from these wires. Internal imperfections are objectionable because they make the rod unsuitable for cold drawing to high strength levels and the extremely fine sizes required. Rod for concrete reinforcement is nondeformed rod produced from steel chemical compositions selected to provide the


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Steel Wire Rod

01 Sep 2005

mechanical property requirements for grade 40 and grade 60, as described in ASTM A 615. This quality rod is produced in coils. Rod for telephone and telegraph wire is produced by practices and to chemical compositions intended for the manufacture of wire having electrical and mechanical properties that will meet the requirements of the various grades of this type of wire.

Special Requirements for Carbon Steel Rod Some of the quality descriptors discussed above imply special requirements for the manufacture and testing of wire rod. A few of the more common requirements are listed below. For some applications, it may be appropriate to add one or more special requirements to those implied by the quality descriptor. Macroetch testing is deep-etch testing to evaluate internal soundness. A representative cross section is etched in a hot acid solution. Fracture Testing. In fracture testing, a specimen is fractured to evaluate soundness and homogeneity. Austenitic Grain Size Requirements. For applications involving carburizing or heat treatment, austenitic grain size for killed steels may be specified as either coarse (grain size 1 through 5) or fine (grain size 5 through 8 inclusive), in accordance with ASTM E 112. Heat-Treating Requirements. When heat-treating requirements must be met in the purchaser's end product, all heat treatment procedures and mechanical property requirements should be clearly specified. Nonmetallic inclusion testing comprises a microscopic examination of longitudinal sections of the rod to determine the nature and frequency of nonmetallic inclusions. Methods B or C of ASTM E 45 are commonly used. Decarburization limits are specified for special applications when required. A specimen is polished so that the entire cross-sectional area is in a single plane, with no rounded edges. After etching with a suitable etchant, the specimen is examined microscopically (usually to 100 diameters), and the results are reported in hundredths of a millimeter or thousandths of an inch. The examination includes the entire periphery, and the results reported should include the amount of free ferrite and the total depth of decarburization. Further details of this microscopic method are contained in SAE Recommended Practice J419, Methods of Measuring Decarburization.

Mechanical Properties of Carbon Steel Rod In the older mills, where rod was coiled hot, there was considerable variation within each coil because of the effect of varying cooling rates from the center to the periphery of the coil. Therefore, as-hot-rolled rod was seldom sold to specific mechanical properties because of the inherent variations of such properties. These properties for a given grade of steel varied from mill to mill and were influenced by both the type of mill and the source of steel being rolled. In new rod mills, which are equipped with controlled cooling facilities, this intracoil variation is kept to a minimum. In such mills, finishing temperature, cooling of water, cooling air, and conveyor speed all are balanced to produce rod with the desired scale and microstructure. This structure, in turn, is reflected in the mechanical properties of the rod and permits the rod to be drawn directly for all but the most demanding applications. The primary source of intracoil variation on these new mills is the overlapping of the coiled rings on the conveyor. These overlapped areas cool at a slower rate than the majority of the ring. Table 1 lists typical values of tensile strength for 5.6 mm (7=32in.) low-carbon steel rod rolled on a modern rod mill equipped with controlled cooling facilities. The values shown are for rods rolled with full air cooling. Tensile strength values for larger-diameter rod are lower, decreasing by approximately 1.9 MPa (270 psi) for each 0.4 mm (1=64in.) increment by which rod diameter exceeds 5.6 mm (7=32in.). Similar analyses of rods rolled without full air cooling or rods rolled on an older mill, where the steel is coiled hot, would be expected to reveal lower tensile strength. Table 1 Tensile strength of 5.6 mm (7=32 in.) diam hot-rolled low-carbon steel rod Data obtained from rod produced with controlled cooling Rimmed

Aluminum killed fine-grain steel

Capped

Silicon killed fine- or coarse-grain steel

Steel grade

MPa

ksi

MPa

ksi

MPa

ksi

MPa

ksi

1005

350

51

...

...

380

55

395

57

1006

360

52

365

53

395

57

405

59

1008

370

54

385

56

405

59

425

62

1010

385

56

400

58

420

61

440

64

1012

405

59

420

61

435

63

455

66

1015

425

62

440

64

450

65

475

69

1017

450

65

455

66

455

66

495

72

1018

...

...

475

69

490

71

525

76

1020

470

68

475

69

485

70

510

74

1022

...

...

...

...

520

75

565

82


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Steel Wire Rod

01 Sep 2005

Table 2 shows typical expected tensile strength values for 5.6mm (7=32in.) medium-high-carbon and high-carbon steel rods rolled on a mill utilizing controlled cooling. The microstructure of such rod approximates that obtained by patenting. The strength generally falls between those obtained by air patenting and lead patenting. Most high-carbon steel wire is drawn from such rods without prior patenting. Tensile strengths for large-diameter rods have been found to average 5.2 MPa (750 psi) lower for each 0.4 mm (1=64in.) increment in diameter over 5.6 mm (7=32in.). Additional strength in control-cooled high-carbon rods can be achieved by using microalloying techniques. These procedures should be used with caution because they may affect other properties such as ductility and durability. Table 2 Tensile strength of 5.6 mm (7=32 in.) diam hot-rolled medium-high-carbon and high-carbon steel rod Data obtained from rod produced with controlled cooling Tensile strength for steel with manganese content of 0.60%

Carbon content of steel, %

MPa

0.30 0.35

0.80% ksi

MPa

641

93

689

100

0.40

745

0.45

793

0.50 0.55

1.00% ksi

MPa

ksi

676

98

717

104

731

106

793

115

108

779

113

820

119

115

834

121

869

126

848

123

883

128

931

135

896

130

938

136

972

141

0.60

951

138

986

143

1020

148

0.65

1000

145

1041

151

1076

156

0.70

1055

153

1089

158

1124

163

0.75

1103

160

1138

165

1179

171

0.80

1151

167

1193

173

1227

178

0.85

1207

175

1241

180

1282

186

Alloy Steel Rod Alloy steels are those steels for which maximum specified manganese content exceeds 1.65% or maximum specified silicon or copper content exceeds 0.60%; or for which a definite range or definite minimum quantity of any other element is specified in order to obtain desired effects on properties. Detailed information on composition ranges and limits of alloy steels can be found in the article "Classification and Designation of Carbon and Low-Alloy Steels" in this Volume.

Qualities and Commodities for Alloy Steel Rod The various qualities of alloy steel wire rod possess characteristics that are adapted to the particular conditions typically encountered during fabrication or service. Manufacture of these steels normally includes careful selection of raw materials for melting, exacting steelmaking practices, selective discard (when the steel is produced in ingots), extensive billet preparation, and extensive testing and inspection. Occasionally, alloy steel of a special quality is specified for the manufacture of wire rod. Aircraft quality alloy steel may be specified for wire rods intended for processing into critical or highly stressed aircraft parts or for similar purposes. Bearing quality alloy steel may be specified for wire rods intended for processing into balls and rollers for antifriction bearings. Bearing quality alloy steel is usually specified when purchasing the standard carburizing grades, such as 4118, 4320, 4620, 4720, and 8620, or the through-hardening, high-carbon chromium grades such as E51100 and E52100. The various standard qualities and commodities available in alloy steel wire rod are described below. Cold heading quality alloy steel rod is used for the manufacture of wire for applications involving cold plastic deformation by such operations as upsetting, heading, forging, or extrusion. Typical parts are fasteners (cap screws, bolts, eyebolts), studs, anchor pins, and balls and rollers for antifriction bearings. Special cold heading quality alloy steel rod is used for wire for applications involving severe cold plastic deformation. Surface quality requirements are more critical than for cold heading quality. Steel with very uniform chemical composition and internal soundness, as well as special surface preparation of the semifinished steel, are required. Typical applications are ball joint suspension studs, socket head screws, recessed-head screws, and valves. Welding quality alloy steel rod is used for the manufacture of wire used as filler metal in electric arc welding or for building up hard wearing surfaces of parts subjected to wear. The heat analysis limits give below for phosphorus and sulfur apply to this quality rod: Maximum percent Steelmaking process


P

S

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Steel Wire Rod

01 Sep 2005

Basic electric

0.025

0.025

Basic open hearth or basic oxygen

0.025

0.035

Special Requirements for Alloy Steel Rod Alloy steel rod can be produced with special requirements in addition to those implied by the quality descriptors discussed above. These special requirements include those given below. Special surface entails a product with minimal frequency and severity of seams and other surface imperfections. Decarburization limits can be specified for special applications. An example of such limits are those shown in the table below for alloy steel rod for wire for heading, forging, roll threading, extrusion, lockwasher, and screwdriver applications. Listed below are the maximum allowable amounts of decarburization as defined by the average value for the depth of the layer of free ferrite plus the layer of partial decarburization (the total affected depth) and the average depth of the layer of free ferrite alone: Average depth of decarburization Free ferrite

Total affected

Nominal rod diameter, mm (in.)

mm

in.

mm

in.

≤6.4 (≤1=4)

0.08

0.003

0.20

0.008

6.8−9.6 (17=64−3=8)

0.08

0.003

0.25

0.010

10−12.8 ( =64− =2)

0.10

0.004

0.30

0.012

25

1

13−18 ( =64− =64)

0.13

0.005

0.36

0.014

18.2−25 ( =32−1)

0.15

0.006

0.41

0.016

33

45

23

When limits closer than those given above are required for the end product, it is sometimes appropriate to incorporate carbon restoration in the fabrication process. For some applications, the rod producer can include carbon restoration in the mill heat treatment. The method of measuring decarburization is the same as that described for carbon steel rods. Heat-Treating Requirements. When the end product must be heat treated, the heat treatment and mechanical properties should be clearly defined. Hardenability requirements are customarily specified by H-steel designations and hardenability bands. These steels and hardenability bands are discussed in the article "Hardenability of Carbon and Low-Alloy Steels" in this Volume. Austenitic Grain Size Determination. Most alloy steels are produced using fine-grain practice. Fine-grain steels are useful in carburized parts, especially when direct quenching is involved, and are less sensitive than coarse-grain steels to variations in heat-treating practices. Coarse-grain steels are deeper hardening and are generally considered more machinable. Austenitic grain size is specified as either coarse (grain sizes 1 through 5) or fine (grain sizes 5 through 8), determined in accordance with ASTM E 112. Nonmetallic-Inclusion Testing. When the nonmetallic-inclusion test is specified, it is commonly done on billets. Prepared and polished specimens are examined microscopically at 100 diameters. Sample locations, number of tests, and limits of acceptability should be established in each instance. Test procedures are described in ASTM E 45. Magnetic-Particle Inspection. For alloy steel rod and wire products subject to magnetic-particle inspection, it is customary for the producer to test the product in a semi-finished form, such as billets (using specimens properly machined from billets), to ensure that the heat conforms to the magnetic-particle inspection requirements, prior to further processing. The method of inspection consists of suitably magnetizing the steel and applying a prepared magnetic powder, either dry or suspended in a suitable liquid, that adheres to the steel along lines of flux leakage. On properly magnetized steel, flux leakage develops along surface or subsurface discontinuities. The results of the inspection will vary with the degree of magnetization, the inspection procedure (including such conditions as relative location of surfaces tested), the method and sequence of magnetizing and applying the powder, and the interpretation. The testing procedure and standards of acceptance for magnetic-particle inspection are described in Aerospace Materials Specification 2301. Macroetch Testing. Soundness and homogeneity of alloy steel rod are sometimes evaluated macroscopically by examining a properly prepared cross section of the product after it has been immersed in a hot acid solution. It is customary to use hydrochloric acid for this purpose. ACKNOWLEDGMENTS The helpful suggestions provided by Zeev Zimerman, Bethlehem Steel Corporation, and Bhaskar Yalamanchili, North Star Steel Texas Company, are greatly appreciated.


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ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Carbon and Low-Alloy Steels Steel Wire Revised by Allan B. Dove, Consultant WIRE can be cold drawn from any of the types of carbon steel or alloy steel rod described in the article "Steel Wire Rod" in this Volume. For convenience, the various grades of carbon steel wire can be divided into the same four classes used for carbon steel rod. Based on carbon content, these classes are: • • • •

Low-carbon steel wire (0.15% C max) Medium-low-carbon steel wire (>0.15 to 0.23% C) Medium-high-carbon steel wire (>0.23 to 0.44% C) High-carbon steel wire (>0.44% C)

The conventional four-digit or five-digit American Iron and Steel Institute⎯Society of Automotive Engineers (AISI-SAE) designation is used to specify the carbon or alloy steel used to make the wire. Carbon and alloy steel wire can be produced in qualities suitable for cold rolling, cold drawing, cold heading, cold upsetting, cold extrusion, cold forging, hot forging, cold coiling, heat treatment, or carburizing and for a wide variety of fabricated products.

Wire Configurations and Sizes Shapes of Wires. Although wire is ordinarily thought of as being only round, it may have any one of an infinite number of sectional shapes, as required by end use. After ordinary round wire, the most common shapes are square, hexagonal, octagonal, oval, half-oval, half-round, triangular, keystone, and flat. In addition to these regular (symmetrical) shapes, wire is also made in various odd and irregular shapes for specific purposes. Flat wire, as defined by AISI, is wire that has been cold rolled or drawn, has a prepared edge, is rectangular in shape, 25 mm (1 in.) or less in width, and less than 9.5 mm (3=8in.) in thickness. Flat wire is generally produced from hot-rolled rods or specially prepared round wire by one or more cold-rolling operations intended primarily for the purpose of obtaining the size and section desired and for improving surface finish, dimensional accuracy, and mechanical properties. Low-carbon steel flat wire can also be produced by slitting cold-rolled flat sheet or strip steel to the desired width. The width-to-thickness ratio and the specified type of edge generally determine the process that will be necessary to produce a specific flat wire item. The edges, finishes, and tempers obtainable in flat wire are similar to those furnished in cold-rolled strip. It should be noted that a product having an approximately rectangular section, rolled from carbon steel round wire of selected size, without edge, is also known as carbon steel flat wire. Sizes of Wire. The size limits for the product commonly known as wire range from approximately 0.13 mm (0.005 in.) to (but not including) 25.4 mm (1 in.) for round sections and from a few tenths of a millimeter to approximately 16 mm (5=8in.) for square sections. Larger rounds and squares (if passed through a die or rolled) and all sizes of hexagonal and octagonal sections are commonly known as cold-drawn bars. The size (diameter) of round wire is expressed in decimal units or by gage numbers. In the United States, the conventional unit is the inch, and wire diameter is determined with micrometers capable of making measurements accurate to at least one thousandth of an inch. Sizes specified are expressed in ten thousandths of an inch, which should be followed by the metric dimension in brackets to two decimal places. There are several different systems of gage numbers that can be used for the measurement of wire, but in general these systems have fallen into disuse and have been replaced by sizes in thousandths of an inch or by metric dimensions. The size of music wire is usually expressed in music wire gage (MWG), which is the standard for this wire application. For iron and steel telephone and telegraph wire, the standard is the Birmingham wire gage (BWG) system. The system commonly used by manufacturers of steel wire (other than the exceptions noted) is the United States steel wire gage (USSWG) or, more commonly, the steel wire gage (SWG) system, and all unidentified gage numbers used in this article will refer to this system. The use of gage numbers for steel wire measurements is falling from favor, and the use of absolute units is gaining acceptance. Table 1 lists decimal equivalents in inches and millimeters for steel wire gage numbers from 7/0 (12.45 mm, or 0.490 in.) to 50 (0.112 mm, or 0.0044 in.). Table 1 Steel wire gage sizes


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Steel Wire

01 Sep 2005

Wire diameter Gage No.

mm

in.

7/0

12.45

0.490

6/0

11.73

0.462

5/0

10.92

0.430

4/0

10.01

0.394

3/0

9.19

0.362

2/0

8.41

0.331

1/0

7.77

0.306

1

7.19

0.283

1 =2

6.91

0.272

2

6.65

0.262

1

6.43

0.253

3

6.20

0.244

31=2

5.94

0.234

1

2 =2

4

5.72

0.225

41=2

5.49

0.216

5

5.26

0.207

51=2

5.08

0.200

6

4.88

0.192

61=2

4.67

0.184

7

4.50

0.177

71=2

4.32

0.170

8

4.11

0.162

8 =2

3.94

0.155

9

3.76

0.148

1

9 =2

3.61

0.142

10

3.43

0.135

10 =2

3.25

0.128

11

3.05

0.120

11 =2

2.87

0.113

12

2.69

0.106

12 =2

2.51

0.099

13

2.34

0.092

1

2.18

0.086

14

2.03

0.080

141=2

1.93

0.076

1

1

1

1

13 =2

15

1.83

0.072

151=2

1.70

0.067

16

1.57

0.062

161=2

1.47

0.058

17

1.37

0.054

171=2

1.30

0.051

18

1.22

0.048

181=2

1.12

0.044

19

1.04

0.041

19 =2

0.97

0.038

20

0.89

0.035

21

0.805

0.0317

1


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Steel Wire

01 Sep 2005

22

0.726

0.0286

23

0.655

0.0258

24

0.584

0.0230

25

0.518

0.0204

26

0.460

0.0181

27

0.439

0.0173

28

0.411

0.0162

29

0.381

0.0150

30

0.356

0.0140

31

0.335

0.0132

32

0.325

0.0128

33

0.300

0.0118

34

0.264

0.0104

35

0.241

0.0095

36

0.229

0.0090

37

0.216

0.0085

38

0.203

0.0080

39

0.191

0.0075

40

0.178

0.0070

41

0.168

0.0066

42

0.157

0.0062

43

0.152

0.0060

44

0.147

0.0058

45

0.140

0.0055

46

0.132

0.0052

47

0.127

0.0050

48

0.122

0.0048

49

0.117

0.0046

50

0.112

0.0044

Wire 20 gage and smaller in size is usually regarded as fine; wire of these sizes is normally drawn and coiled on 203 mm (8 in.) diam blocks. Larger blocks are used as finished wire diameter increases. For example, 2.34 or 0.092 in. (13 gage) wire is generally drawn on 559 mm (22 in.) blocks. Table 2 indicates the usual block sizes by gages for wires between 0.889 and 12.70 mm (0.035 and 0.500 in.). Table 2 Wiredrawing block sizes and corresponding coil diameters for coarse round wire Coil weight 0.55% C) are unsuitable because they are notch sensitive. Fig. 9 Fatigue data for 1040 and 4037 steel bolts. The bolts ( 3=8by 2 in., 16 threads to the inch) had a hardness of 35 HRC.


Page 495


Threaded Steel Fasteners

01 Sep 2005

Tensile properties of the 1040 steel at three-thread exposure were yield strength, 1060 MPa (154 ksi); tensile strength (axial), 1200 MPa (175 ksi); tensile strength (wedge), 1190 MPa (173 ksi). For the 4037 steel: yield strength, 1110 MPa (161 ksi); tensile strength (axial), 1250 MPa (182 ksi); tensile strength (wedge), 1250 MPa (182 ksi)

The principal design feature of a bolt is the threaded section, which establishes a notch pattern inherent in the part because of its design. The form of the threads, plus any mechanical or metallurgical condition that also creates a surface notch, is much more important than steel composition in determining the fatigue strength of a particular lot of bolts. Some of these factors are discussed below. Causes and Prevention of Fatigue Crack Initiation. The origin of a fatigue crack is usually at some point of stress concentration, such as an abrupt change of section, a deep scratch, a notch, a nick, a fold, a large inclusion, or a marked change in grain size. Fatigue failures in bolts often occur in the threaded section immediately adjacent to the edge of the nut (or mating part) on the washer side, at or near the first thread inside the nut (or mating part). This area of stress concentration occurs because the bolt elongates as the nut is tightened, thus producing increased loads on the threads nearest the bearing face of the nut, which add to normal service stresses. This condition is alleviated to some extent by using nuts of a softer material that will yield and distribute the load more uniformly over the engaged threads. Significant additional improvement in fatigue life is also obtained by rolling (cold working) the threads rather than cutting them and also by rolling threads after heat treatment rather than before. Other locations of possible fatigue failure of a bolt under tensile loading are the thread runout and the head-to-shank fillet. Like the section of the bolt thread described in the previous paragraph, these two locations are also areas of stress concentration. Any measures that decrease stress concentration can lead to improved fatigue life. Typical examples of such measures are the use of UNJ increased root radius threads (see MIL-S-8879A) and the use of internal thread designs that distribute the load uniformly over a large number of bolt threads. Shape and size of the head-to-shank fillet are important, as is a generous radius from the thread runout to the shank. In general, the radius of this fillet should be as large as possible while at the same time permitting adequate head-bearing area. This requires a design trade-off between the head-to-shank radius and the head-bearing area to achieve optimum results. Cold working of the head fillet is another common method of preventing fatigue failure because it induces a residual compressive stress and increases the material strength. Several other factors are also important in avoiding fatigue fracture at the head-to-shank fillet. The heads of most fasteners are formed by hot or cold forging, depending on the type of material and size of the bolt. In addition to being a relatively low-cost manufacturing method, forging provides smooth, unbroken grain flow lines through the head-to-shank fillet, which closely follows the external contour of the bolt (Fig. 10 ) and therefore minimizes stress raisers, which promote fatigue cracking. In the hot forging of fastener heads, temperatures must be carefully controlled to avoid overheating, which may cause grain growth. Several failures of 25 mm (1 in.) diam type H-11 airplane-wing bolts quenched and tempered to a tensile strength of 1800 to 1930 MPa (260 to 280 ksi) have been attributed to stress concentration that resulted from a large grain size in the shank. Other failures in these 25 mm (1 in.) diam bolts, as well as in other similarly quenched and tempered steel bolts, were the result of cracks in untempered martensite that formed as a result of overheating during finish grinding. Fig. 10 Uniform, unbroken grain flow around the contours of the forged head of a threaded fastener. The uniform, unbroken grain flow minimizes stress raisers and unfavorable shear planes and therefore improves fatigue strength.


Page 496



01 Sep 2005

Influence of the Thread-Forming Method on Bolt Fatigue Strength. The method of forming the thread is an important factor influencing the fatigue strength of bolts. Specifically, there is a marked improvement when threads are rolled rather than either cut or ground, particularly when the threads are rolled after the bolt has been heat treated (Fig. 11 ). Fig. 11 Fatigue limits for roll-threaded steel bolts. (a) Four different lots of bolts that were roll threaded, then heat treated to average hardness of 22.7, 26.6, 27.6, and 32.6 HRC. (b) Five different lots that were heat treated to average hardnesses of 23.3, 27.4, 29.6, 31.7, and 33.0 HRC, then roll threaded. Bolts having higher hardnesses in each category had higher fatigue strengths.

Other factors being equal, a bolt with threads properly rolled after heat treatment⎯that is, free from mechanical imperfections⎯has a higher fatigue limit than one with cut threads. This is true for any strength category. The cold work of rolling increases the strength at the weakest section (the thread root) and imparts residual compressive stresses, similar to those imparted by shot peening. The larger and smoother root radius of the rolled thread also contributes to its superiority. Because of the fatigue life concern, all bolts and screws greater than grade 1 and less than 19 mm (3=4in.) in diameter and 150 mm (6 in.) in length are to be roll threaded. Studs and larger bolts and screws may have the threads rolled, cut or ground. Effect of Surface Treatment on Fatigue. Light cases, such as from carburizing or carbonitriding, are rarely recommended and should not be used for critical externally threaded fasteners, such as bolts, studs, or U-bolts. The cases are quite brittle and crack when the fasteners are tightened or bent in assembly or service. These cracks may then lead to fatigue cracking and possible fracture. Chromium and nickel platings decrease the fatigue strength of threaded sections and should not be used except in a few applications, such as automobile bumper studs or similar fasteners that operate under conditions of low stress and require platings for appearance. Cadmium and zinc platings slightly reduce fatigue strength. Electroplated parts for high-strength applications should be treated after plating to eliminate or minimize hydrogen embrittlement (which is a strong contributor to fatigue cracking). Installation. As noted at the beginning of this section on fatigue failures, bolt loading is a major factor in the fatigue failure of threaded fasteners. When placed into service, bolts are most likely to fail by fatigue if the assemblies involve soft gaskets or flanges, or if the bolts are not properly aligned and tightened. Fatigue resistance is also related to clamping force. In many assemblies, a certain minimum clamping force is required to ensure both proper alignment of the bolt in relation to other


Page 497



01 Sep 2005

components of the assembly and proper preload on the bolt. The former ensures that the bolt will not be subjected to undue eccentric loading, and the latter that the correct mean stress is established for the application. In some cases, clamping stresses that exceed the yield strength may be desirable; experiments have shown that bolts clamped beyond the yield point have better fatigue resistance than bolts clamped below the yield point. Decreasing the bolt stiffness can also reduce cyclic stresses. Methods commonly used are reduction of the cross-sectional area of the shank to form a waisted shank and rolling the threads further up the exposed shank to increase the "spring" length. Stress-corrosion cracking (SCC) is an intergranular fracture mechanism that sometimes occurs in highly stressed fasteners after a period of time, and it is caused by a corrosive environment in conjunction with a sustained tensile stress above a threshold value. An adverse grain orientation increases the susceptibility of some materials to stress corrosion. Consequently, SCC can be prevented by excluding the corrodent, by keeping the static tensile stress of the fastener below the critical level for the material and grain orientation involved, or by changing to a less susceptible material or material condition. Because tensile loads (even residual tensile loads) are required to produce SCC, compressive residual stresses may prevent SCC. As with the environmentally induced cracking from hydrogen embrittlement and liquid-metal embrittlement (see the article "Embrittlement of Steels" in this Volume), the understanding of SCC is largely phenomenological, without any satisfactory mechanistic model for predicting SCC or the other forms of environmentally induced cracking. This lack of mechanistic predictability of SCC is particularly unfortunate because measurable corrosion usually does not occur before or during crack initiation or propagation. Even when corrosion does occur, it is highly localized (that is, pitting, crevice attack) and may be difficult to detect. Moreover, SCC is a complex synergistic phenomenon resulting from the combined simultaneous interaction of mechanical and chemical conditions. Pre-corrosion followed by loading in an inert environment will not result in any observable crack propagation, while simultaneous environmental exposure and application of stress will cause time-dependent subcritical crack propagation. The susceptibility of a metal to SCC depends on the alloy and the corrodent. The National Association of Corrosion Engineers, the Materials Technology Institute of the Chemical Process Industries, and others have published tables of corrodents known to cause SCC of various metal alloy systems (Ref 2, 3, 4). This literature should be used only as a guide for screening candidate materials for further in-depth investigation, testing, and evaluation. In general, plain carbon steels are susceptible to SCC by several corrodents of economic importance, including aqueous solutions of amines, carbonates, acidified cyanides, hydroxides, nitrates, and anhydrous ammonia. Carbon steels, low-alloy steels, and H-11 tool steels with ultimate tensile strengths above 1380 MPa (200 ksi) are susceptible to SCC in a seacoast environment. Of the various bolt steels, bolts made from ISO class 12.8 have experienced failures for SCC in automotive applications (Ref 1). Stainless steels are also susceptible to SCC in some environments. Even though the micromechanistic causes of SCC are not entirely understood, some investigators consider SCC to be related to hydrogen damage and not strictly an active-path corrosion phenomenon. Although hydrogen can be a factor in the SCC of certain alloys (see Example 1 ), sufficient data are not available to generalize this concept. For example, SCC can be assisted by such factors as nuclear irradiation. More information on SCC is available in Corrosion, Volume 13 of ASM Handbook. Example 1: Hydrogen-Assisted SCC Failure of Four AISI 4137 Steel Bolts. Figure 12 shows an example of hydrogen-assisted SCC failure of four AISI 4137 steel bolts having a hardness of 42 HRC. Although the normal service temperature (400 °C, or 750 °F) was too high for hydrogen embrittlement, the bolts were also subjected to extended shutdown periods at ambient temperatures. The corrosive environment contained trace hydrogen chloride and acetic acid vapors as well as calcium chloride if leaks occurred. The exact service life was unknown. The bolt surfaces showed extensive corrosion deposits. Cracks had initiated at both the thread roots and the fillet under the bolt head. Figure 12 (b) shows a longitudinal section through the failed end of one bolt. Multiple, branched cracking was present, typical of hydrogen-assisted SCC in hardened steels. Chlorides were detected within the cracks and on the fracture surface. The failed bolts were replaced with 17-4 PH stainless steel bolts (Condition H 1150M) having a hardness of 22 HRC (Ref 5). Fig. 12 4137 steel bolts (hardness: 42 HRC) that failed by hydrogen-assisted SCC caused by acidic chlorides from a leaking polymer solution. (a) Overall view of failed bolts. (b) Longitudinal section through one of the failed bolts in (a) showing multiple, branched hydrogen-assisted stress-corrosion cracks initiating from the thread roots. Source: Ref 5


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01 Sep 2005

Fabrication Most bolts are made by cold heading. Other cold-forming methods, including cold extrusion, are used in the manufacture of threaded fasteners. Current technology is such that not only low-carbon steels but also medium-carbon and even low-alloy steels can be successfully impact extruded. Parts having heads that are large in relation to the shank diameter can be hot headed or produced cold on a two-die, three-blow machine. Hot heading is also more practical for bolts with diameters larger than about 32 mm (11=4in.) because of equipment limitations and increased probability of tool failures with cold heading. Platings and Coatings. Mots carbon steel fasteners are plated or coated. Common coatings are zinc, cadmium, and phosphate and oil. Other supplementary finishes are gaining popularity, especially in critical applications. The principal reason for fastener plating or coating is corrosion resistance (see the section "Corrosion Protection" in this article), although the appearance and the installation torque-tension relationship are also improved. Plating is the deposition of metal onto the surface of the base metal. For commercial applications, plating is achieved by electroplating, hot dipping, or mechanical application. In general, the addition of plating increases the dimensions of the fastener by two times the plating thickness and by almost four times the plating thickness in the thread dimensions. The thread assembly may be affected by the increase in fastener size due to plating. High-strength fasteners, usually with higher carbon content, are susceptible to hydrogen embrittlement when being acid cleaned, electrocleaned, or electroplated. Hydrogen penetration into a fastener can be minimized by baking at about 190 to 200 °C (375 to 400 °F) for 3 to 24 h. For applications in which hydrogen embrittlement is a concern or for a critical application of high-strength fasteners, the mechanical application of plating should be considered.

Clamping Forces To operate effectively and economically, threaded fasteners should be designed to be torqued near the proof stress, as dictated by the cross-sectional area of the load-carrying parts of the fastener and the desired clamping force. The actual clamping force attained in any assembly will be influenced by such factors as: • Roughness of the mating surfaces • Coatings, contaminants, or lubricants on the mating surfaces • Platings or lubricants on the threads Typically, torque values are established to result in a clamping force equal to about 75% of the proof load. For some applications, bolts are torqued beyond their proof stress with no detrimental results, provided they are permanent fasteners holding static loads. Because it is difficult to measure bolt tension (clamping force) in production installations, torque values are used in most applications. Some critical joints in assembly-line processes are using torque-angle or torque-to-yield methods of tightening. In the torque-angle method, the bolt is torqued to a low seating torque level to mate all surfaces, then rotated a specific angle. The angle rotation has a linear relationship with extension because of the constant pitch and therefore with clamp load. In the torque-to-yield method, torque and angular rotation are monitored during, installation by a microprocessor and bolt rotation continues until the relationship between the two is not linear. This point is defined as the yield point in torque tension. The clamping forces generated at given torques are very dependent on the coefficient of friction at the threads and at the bearing face; therefore, they are highly dependent on fastener coatings. Common fastener coatings are zinc, cadmium, and phosphate and oil. The maximum clamping force that can be effectively employed in any bolt is often limited by the compressive strength of the materials being bolted. If this value is exceeded, the bolt head or nut will be pulled into the parts being bolted, with a subsequent reduction in clamping force. The assembly then becomes loose, and the bolt is susceptible to fatigue failure. If high-tensile bolts are necessary to join low compressive strength materials, hardened washers should be used under the head of the bolt and under the nut to distribute bearing pressure more evenly and to avoid the condition described above. The value of high clamping forces, apart from lessening the possibility of the nut loosening, is that the working stresses (against solid abutments) are always less than the clamping forces induces in a properly selected bolt. This ensures against cyclic stress and possible fatigue failure. REFERENCES 1. T.J. Hughel, "Delayed Fracture of Class 12.8 Bolts in Automotive Rear Suspensions," SAE Technical Paper Series 820122, Society of Automotive Engineers, 1982 2. Corrosion Data Survey⎯Metals Section , 5th ed., National Association of Corrosion Engineers, 1974, p. 268−269 3. D.R. McIntyre and C.P. Dillon, Guidelines for Preventing Stress Corrosion Cracking in the Chemical Process Industries , Publication 15, Materials Technology Institute of the Chemical Process Industries, 1985, p 8−14 4. The Role of Stainless Steels in Petroleum Refining, American Iron and Steel Institute, 1977, p 41 5. D. Warren, Hydrogen Effects on Steel, in Process Industries Corrosion, National Association of Corrosion Engineers, 1986,


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p 31−43


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ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Carbon and Low-Alloy Steels Steel Springs Revised by Loren Godfrey, Associated Spring/Barnes Group, Inc. STEEL SPRINGS are made in many types, shapes, and sizes, ranging from delicate hairsprings for instrument meters to massive buffer springs for railroad equipment. The major portion of this article discusses those relatively small steel springs that are cold wound from wire. Relatively large, hot-wound springs are quite different from cold-wound springs and are treated in a separate section. Flat and leaf springs are also treated separately to the extent that they differ from wire springs in material and fabrication. Wire springs are of four types: compression springs (including die springs), extension springs, torsion springs, and wire forms. Compression springs are open wound with varying space between the coils and are provided with plain, plain and ground, squared, or squared and ground ends. The spring can be cylindrical, conical, barrel, or hour glass shaped. Extension springs are normally close wound, usually with specified initial tension, and, because they are used to resist pulling forces, are provided with hook or loop ends to fit the specific application. Ends may be integral parts of the spring or specially inserted forms. Torsion springs are usually designed to work over an arbor and to resist a force that causes the spring to wind. Wire forms are made in a wide variety of shapes and sizes. Flat springs are usually made by stamping and forming of strip material into shapes such as spring washers. However, there are other types, including motor springs (clock type), constant-force springs, and volute springs, that are wound from strip or flat wire. Chemical composition, mechanical properties, surface quality, availability, and cost are the principal factors to be considered in selecting steel for springs. Both carbon and alloy steels are used extensively. Steels for cold-wound springs differ from other constructional steels in four ways. • • • •

They are cold worked more extensively They are higher in carbon content They can be furnished in the pretempered condition They have higher surface quality

The first three items increase the strength of the steel, and the last improves fatigue properties. For flat, cold-formed springs made from steel strip or flat wire, narrower ranges of carbon and manganese are specified than for cold-wound springs made from round or square wire. Where special properties are required, spring wire or strip made of stainless steel, a heat-resistant alloy, or a nonferrous alloy can be substituted for the carbon or alloy steel, provided that the design of the spring is changed to compensate for the differences in properties between the materials (see the section "Design" in this article). Table 1 lists grade, specification, chemical composition, properties, method of manufacture, and chief applications of the materials commonly used for cold-formed springs. Hot-formed carbon and alloy steel springs are discussed in this article. Table 1 Common wire and strip materials used for cold-formed springs Tensile properties

Material type

Nominal Grade and compositio specification n, %

Minimum tensile strength(a), MPa (ksi)

Modul us of elastici ty (E), GPa (psi × 106)

1590−2750 (230−399)

210 (30)

Torsion properties Design stress, % of Modulus of Max minim rigidity allowable um (G), tensile Hardnes temperature , strengt GPa (psi s, × 106) °C ( °F) h(b) HRC(c)

Method of manufacture, chief applications, special properties

Cold drawn wire High-carbo Music wire, n steel ASTM A

C 0.70−1.00,


45

80 (11.5)

41−60

120 (250)

Drawn to high and uniform tensile strength;

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228

Steel Springs

01 Sep 2005

for high-quality springs and wire forms

Mn 0.20−0.60

Hard drawn, C ASTM A 0.45−0.85, 227 Mn 0.30−1.30

Class I 1010−1950 (147−283) Class II 1180−2230 (171−324)

210 (30)

40

80 (11.5)

31−52

120 (250)

For average-stress applications; lower-cost springs and wire forms

C 0.65−1.00, Mn 0.20−1.30

1640−2410 (238−350)

210 (30)

45

80 (11.5)

41−60

120 (250)

For higher-quality springs and wire forms

Oil tempered, C ASTM A 0.55−0.85, 229 Mn 0.30−1.20

Class I 1140−2020 (165−294) Class II 1320−2330 (191−324)

210 (30)

45

80 (11.5)

42−55

120 (250)

Heat treated before fabrication; for general-purpose springs

C 0.60−0.75, Mn 0.60−0.90

1480−1650 (215−240)

210 (30)

45

80 (11.5)

45−49

120 (250)

Heat treated before fabrication; good surface condition and uniform tensile strength

Alloy steel Chromium vanadium, ASTM A 231, A 232(d)

C 0.48−0.53, Cr 0.80−1.10, V 0.15 min

1310−2070 (190−300)

210 (30)

45

80 (11.5)

41−55

220 (425)

Heat treated before fabrication; for shock loads and moderately elevated temperature; ASTM A 232 for valve springs

Modified chromium vanadium VSQ(d), ASTM A 878

C 0.60−0.75, Cr 0.35−0.60, V 0.10−0.25

1410−2000 (205−290)

...

...

...

...

...

Heat treated before fabrication; for valve springs and moderately elevated temperatures

Chromium silicon, ASTM A 877(d), A 401

C 0.51−0.59, Cr 0.60−0.80, Si 1.20−1.60

1620−2070 (235−300)

210 (30)

45

80 (11.5)

48−55

245 (475)

Heat treated before fabrication; for shock loads and moderately elevated temperature; ASTM A 877 for valve springs

Type 302(18−8), ASTM A 313

Cr 17−19, Ni 8−10

860−2240 (125−325)

190 (28)

30−40 69 (10.0)

35−45

290 (550)

General-purpose corrosion and heat resistance; magnetic in spring temper

Type 316, ASTM A 313

Cr 16−18, Ni 10−14, Mo 2−3

760−1690 (110−245)

190 (28)

40

69 (10.0)

35−45

290 (550)

Good heat resistance; greater corrosion resistance than 302; magnetic in spring temper

Type 631 (17−7 PH), ASTM A 313

Cr 16−18, Ni 6.50−7.75, Al 0.75−1.50

Condition CH-900 1620−2310 (235−335)

200 (29.5)

45

76 (11.0)

38−57

340 (650)

Precipitation hardened after fabrication; high strength and general-purpose corrosion resistance; magnetic in spring temper

Nonferrous Copper alloy Cu 94−96, alloys 510 Sn 4.2−5.8 (phosphor bronze A), ASTM B 159

720−1000 (105−145)

100 (15)

40

43 (6.25) 98−104(e )

90 (200)

Good corrosion resistance and electrical conductivity

High-tensile hard drawn, ASTM A 679

Carbon VSQ(d), ASTM A 230

Stainless steel


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01 Sep 2005

Copper alloy Cu 98, Be 170 1.8−2.0 (beryllium copper) ASTM B 197

1030−1590 (150−230)

130 (18.5)

45

50 (7.0)

35−42

200 (400)

Can be mill hardened before fabrication; good corrosion resistance and electrical conductivity; high mechanical properties

Monel 400, AMS 7233

1000−1240 (145−180)

180 (26)

40

65 (9.5)

23−32

230 (450)

Good corrosion resistance at moderately elevated temperature

110−1380 (160−200)

180 (26)

40

65 (9.5)

23−35

290 (550)

Excellent corrosion resistance at moderately elevated temperature

1100−1380 (160−200)

200 (29)

35

72 (10.4)

35−42

510 (950)

Precipitation hardened after fabrication; good corrosion resistance at elevated temperature

Inconel 600, Ni 76, Cr QQ-W-390(f 15.8, Fe 7.2 )

1170−1590 (170−230)

215 (31)

40

76 (11.0)

35−45

370 (700)

Good corrosion resistance at elevated temperature

Inconel 718

1450−1720 (210−250)

200 (29)

40

77 (11.2)

45−50

590 (1100) Precipitation hardened after fabrication; good corrosion resistance at elevated temperature

No. 1 temper 1070 (155) Spring temper 1310−1590 (190−230

215 (31)

40

83 (12.0)

No. 1 34−39 Spring 42-48

400−600 Precipitation hardened (750−1100) after fabrication; good corrosion resistance at elevated temperature

Ni 66, Cu 31.5

Ni 65, Cu Monel 29.5, Al 2.8 K-500, QQ-N-286(f) High-tempe A-286 alloy rature alloys

Fe 53, Ni 26, Cr 15

Ni 52.5, Cr 18.6, Fe 18.5

Ni 73, Cr Inconel 15, Fe 6.75 X-750, ;AMS 5698, 5699 Cold-rolled strip Carbon steel

Medium carbon (1050), ASTM A 682

C 0.47−0.55, Mn 0.60−0.90

Tempered 1100−1930 (160−280)

210 (30)

...

...

Annealed 85 max(e), tempered 38−50

120 (250)

General-purpose applications

"Regular" carbon (1074), ASTM A 682

C 0.69−0.80, Mn 0.50−0.80

Tempered 1100−2210 (160−320)

210 (30)

...

...


120 (250)

Most popular material for flat springs

High carbon (1095), ASTM A 682

C 0.90−1.04, Mn 0.30-0.50

Tempered 1240−2340 (180−340)

210 (30)

...

...


120 (250)

High-stress flat springs

C 0.48−0.53, Cr 0.80-0.10, V 0.15 min

1380−1720 (200−250)

210 (30)

...

...

42−48

220 (425)

Heat treated after fabrication; for shock loads and moderately elevated temperature

C Chromium silicon, AISI 0.51−0.59, 9254 Cr 0.60−0.80, Si 1.20−1.60

1720−2240 (250−325)

210 (30)

...

...

47−51

245 (475)

Heat treated after fabrication; for shock loads and moderately elevated temperature

Type 301

Cr 16−18, Ni 6−8

1655−2650 (240−270)

190 (28)

...

...

48−52

150 (300)

Rolled to high yield strength; magnetic in spring temper

Type 302 (18−8)

Cr 17−19, Ni 8−10

1280−1590 (185−230)

190 (28)

...

...

42−48

290 (550)

General-purpose corrosion and heat resistance; magnetic in spring temper

Alloy steel Chromium vanadium, AMS 6455

Stainless steel


Page 503


Steel Springs

01 Sep 2005

Type 316

Cr 16−18, Ni 10−14, Mo 2−3

1170−1590 (170−230)

190 (28)

...

...

38−48

290 (550)

Good heat resistance; greater corrosion resistance than 302; magnetic in spring temper

Type 631 (17−7 PH), ASTM A 693

Cr 16−18, Ni 6.50−7.75, Al 0.75−1.50

Condition CH-900 1655 (240)

200 (29)

...

...

46 min

340 (650)

Precipitation hardened after fabrication; high strength and general-purpose corrosion resistance; magnetic in spring temper

650−750 (95−110)

100 (15)

...

...

94−98(e)

90 (200)

Good corrosion resistance and electrical conductivity

Copper alloy Cu 98, Be 170 1.6−1.8 (beryllium copper), ASTM B 194

1240−1380 (180−200)

130 (18.5)

...

...

39 min

200 (400)

Can be mill hardened before fabrication; good corrosion resistance and electrical conductivity; high mechanical properties

Monel 400, AMS 4544

690−970 (100−140)

180 (26)

...

...

98 min(e)

230 (450)

Good corrosion resistance at moderately elevated temperature

1170−1380 (170−200)

180 (26)

...

...

34 min

290 (550)

Excellent corrosion resistance at moderately elevated temperature

1100−1380 (160−200)

200 (29)

...

...

30−40

510 (950)

Precipitation hardened after fabrication; good corrosion resistance at elevated temperature

Inconel 600, Ni 76, Cr 15.8, Fe 7.2 ASTM B 168, AMS 5540

1000−1170 (145−170)

215 (31)

...

...

30 min

370 (700)

Good corrosion resistance at elevated temperature

Inconel 718, Ni 52.5, Cr AMS 5596, 18.6, Fe 18.5 AMS 5597

1240−1410 (180−204)

200 (29)

...

...

36

Ni 73, Cr Inconel X-750, AMS 15, Fe 6.75 5542

1030 (150)

215 (31)

...

...

30 min

Nonferrous Copper alloy Cu 94−96, alloys 510 Sn 4.2−5.8 (phosphor bronze A), ASTM B103

Ni 66, Cu 31.5

Monel K-500 Ni 65, Cu QQ-N-286(f) 29.5, Al 2.8 High-tempe A-286 alloy, Fe 53, Ni AMS 5525 26, Cr 15 rature alloys

590 (1100) Precipitation hardened after fabrication; good corrosion resistance at elevated temperature

400−590 Precipitation hardened (750−1100) after fabrication; good corrosion resistance at elevated temperature (a) Minimum tensile strength varies within the given range according to wire diameter (see the applicable specification). Maximum tensile strength is generally about 200 MPa (30 ksi) above the minimum tensile strength. (b) For helical compression or extension springs; design stress of torsion and flat springs taken as 75% of minimum tensile strength. (c) Correlation between hardness and tensile properties of wire is approximate only and should not be used for acceptance or rejection. (d) Valve-spring quality. (e) HRB values. (f) Federal specification. Source: Ref 1

Mechanical Properties Steels of the same chemical composition may perform differently because of different mechanical and metallurgical characteristics. These properties are developed by the steel producer through cold work and heat treatment or by the spring manufacturer through heat treatment. Selection of round wire for cold-wound springs is based on minimum tensile strength for each wire size and grade (Fig. 1 ) and on minimum reduction in area (40% for all sizes). Fig. 1 Minimum tensile strength of steel spring wire. VSQ, valve-spring quality


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Steel Springs

01 Sep 2005

Rockwell hardness and tensile strength for any grade of spring steel strip depend on thickness. The same properties in different thicknesses can be obtained by specifying different carbon contents. The relationship between thickness of spring steel strip containing 0.50 to 0.95% C and Rockwell hardness is shown in Fig. 2 . The optimum hardness of a spring steel increases gradually with decreasing thickness. Fig. 2 Effect of strip thickness on the optimum hardness of spring steel strip for high-stress use. Hardness on HRC scale may be lowered 3 to 4 points for greater toughness. Instability of ductility is sometimes encountered above 57 HRC.

The hardness scale that can be used for thin metal depends on the hardness and the thickness of the metal (see Table 3 in the article "Rockwell Hardness Testing" in Mechanical Testing, Volume 8 of ASM Handbook, formerly 9th Edition Metals Handbook). For testing spring steel strip, which has a minimum hardness of 38 HRC, the Rockwell C scale is used for metal thicker than 0.89 mm (0.035 in.). For thickness ranges of 0.89 to 0.64 mm (0.035 to 0.025 in.), 0.64 to 0.5 mm (0.025 to 0.020 in.), and 0.5 to 0.33 mm (0.020 to 0.013 in.), the Rockwell 45N, 30N, and 15N scales are preferred. For thickness less than 0.33 mm (0.013 in.), the Vickers (diamond pyramid) scale is recommended. As the strip hardness increases, the thickness that can be safely tested decreases. It has been found that the readings obtained with the Vickers indentor are less subject to variation in industrial circumstances than those obtained with the Knoop indentor. The 500 g load Vickers test is used for spring steel strip in thicknesses as low as 0.08 mm (0.003 in.). If readings are made using the proper hardness scale for a given thickness and hardness, they can be converted to HRC values using charts like those in the appendix to the article "Miscellaneous Hardness Tests" in Mechanical Testing, Volume 8 of ASM Handbook, formerly 9th Edition Metals Handbook. Similar charts appear in ASTM A 370 and in the cold-rolled flat wire section of the Steel Products Manual of the American Iron and Steel Institute (AISI). Chart No. 60 published by Wilson Instrument Division, American Chain & Cable Company, Inc., can also be used for this conversion. For specific steel springs, hardness can be held to within 4 points on the Rockwell C scale.


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01 Sep 2005

Note that in Table 1 and in the section "Design" in this article, design-stress values are given as percentages of minimum tensile strength. These values apply to springs that are coiled or formed and then stress relieved, which are used in applications involving relatively few load cycles. If each spring is coiled or formed so as to allow for some set, and then deflected beyond the design requirements, higher design stresses can be used. This is discussed in the section "Residual Stresses" in this article. As a further aid in selecting steels for springs, Table 2 lists the suitable choices for cold-wound helical springs in various combinations combinations of size, stress, and service. Each recommendation is the most economical steel that will perform satisfactorily under the designated conditions and that is commercially available in the specific size. Table 2 Recommended ASTM grades of carbon and alloy steel wire for cold-wound helical springs See Table 1 for the type of wire and the composition of the ASTM grades given below. Corrected maximum working stress(a) Diameter of spring wire(b)

MPa

ksi

0.31−0.51 0.89−3.18 mm mm (0.005−0.020 0.51−0.89 mm (0.035−0.125 (0.020−0.035 in.) in.) in.)

3.18−6.35 mm (0.125−0.250 in.)

6.35−12.70 mm (0.250−0.500 in.)

1270−15.88 mm (0.500−0.625 in.)

Compression springs, static load (set removed, springs stress relieved)(c) 550

80

A 228(d)

A 227(d)

A 227(d)

A 227

A 227

227(e), A 229

690

100

A 228(d)

A 227(d)

A 227(d)

A 227

A 227(f), A 229

A 229

825

120

A 228(d)

A 227(d)

A 227

227(g), A 229

A 229(h)

...

965

140

A 228

A 227

A 227(i), A 229 A 229

A 401(h)

...

1100

160

A 228

A 227

A 228, A 229

A 229(j), A 228

A 401(k)(h)

...

1240

180

A 228

A 228

A 228

A 228(l)(h)

...

...

1380

200

A 228

A 228

A 228(m)(h)

...

...

...

1515

220

A 228

A 228(n)(h)

...

...

...

...

1655

240

A 228(o)(h)

...

...

...

...

...

Compression springs, variable load, designed for minimum life of 100,000 cycles (set removed, springs stress relieved)(p) 550

80

A 228(d)

A 227(d)

A 227(d)

A 227(q), A 229

A 229(r), A 401

A 401

690

100

A 228(d)

825

120

A 228(d)

A 227(d)

A 227, A 229

A 229(s), A 401

A 401

A 401

A 227

A 227(t), A 229 A 229(u), A 401

965

140

A 228

A 229

A 229(w), A 228

1100

160

A 228

A 228

A 228(w)(h), A 401

1240

180

A 228

A 228(y)(h)

1380

200

A 228(z)(h)

...

A 401(v)

A 228(x)

... ...

...

...

...

...

...

...

...

...

...

...

...

...

Compression springs, dynamic load, designed for minimum life of 10 million cycles (set removed, springs stress relieved)(p) 415

60

A 228(d)

A 227(d)

A 227(d)

550

80

A 228(d)

A 228(d)

A 229(t), A 228 A 230

A 227(u), A 230

690

100

A 228(d)

A 228

A 228(aa), A 230

825

120

A 228(d)

A 228

A 230(cc)

A 229(k)(h), A 230

...

A 230

...

A 230(bb)(h) ...

...

...

...

...

Compression and extension springs, static load (set not removed, compression springs stress relieved)(c) 550

80

A 228

A 227

A 227

690

100

A 228

A 227

A 227(t), A 229 A 401

A 227

A 227(dd), A 229 A 401(h)

A 229 ...

825

120

A 228

A 227

A 227(ee), A 229

A 401

A 401(k)(h)

...

965

140

A 228

A 228

A 228(ff), A 401

A 401(h)

...

...

1100

160

A 228(h)

...

...

...

...

...

1240

180

A 228(o)(h)

...

...

...

...

...

Compression and extension springs, designed for minimum life of 100,000 cycles (set not removed, compression springs stress relieved)(p) 415

60

A 228

A 227


A 227

A 227

A 227

A 229

Page 506


Steel Springs

01 Sep 2005

550

80

A 228

A 227

A 227

A 227(x), A 229

A 229

690

100

A 228

A 229

A 229(cc), A 228

A 228(s), A 401

A 401(h)

825

120

A 228

A 228

A 228(w)

A 401(h)

965

140

A 228

A 228(h)

1100

160

A 228(o)(h)

...

A 401 ... ...

...

...

...

...

...

...

...

...

...

Compression and extension springs, designed for minimum life of 10 million cycles (set not removed, compression springs stress relieved)(p) 275

40

A 228

A 227

A 227

A 227

415

60

A 228

A 227

A 227(aa), A 230

A 230(bb)(h)

550

80

A 228

A 228

A 228(aa)(h), A 230

690

100

A 228(o)

...

...

A 227(m), A 229

A 229

...

...

...

...

...

...

...

...

Torsion springs (springs not stress relieved)(p) 690

100

A 228

A 227

A 227

A 227(bb), A 229

A 229(v), A 401

...

825

120

A 228

A 227

A 227(w), A 229

A 229(bb), A 228

A 401(r)(h)

...

965

140

A 228

A 229

A 229(i), A 228 A 228(q)(h)

...

...

1100

160

A 228

A 228

A 228(w)(h)

...

...

...

1240

180

A 228

A 228(h)

...

...

...

...

1380 200 A 228(h) ... ... ... ... ... (a) Stress corrected by the Wahl factor. See the section "Wahl Correction" in this article. (b) To 1.37 mm (0.054 in.) (c) Where more than one steel is shown for an indicated range of wire diameter, the first is recommended up to the specific diameter listed in the footnote referred to; the last steel listed in any multiple choice is recommended for the remainder of the indicated wire diameter range. (d) To 5.26 mm (0.207 in.) (e) Shot peening is not necessary for statically loaded springs. (f) To 2.03 mm (0.080 in.) (g) Set removal not required in this range. (h) To 7.19 mm (0.283 in.). (i) To 14.29 mm (0.563 in.). (j) To 1.12 mm (0.044 in.). (k) To 10.32 mm (0.406 in.). (l) To 2.69 mm (0.105 in.) (m) To 4.11 mm (0.162 in.). (n) Yielding likely to occur beyond this limit. (o) To 1.83 mm (0.072 in.) (p) To 3.81 mm (0.150 in.). (q) To 11.11 mm 0.437 in.) (r) To 5.33 mm (0.210 in.). (s) To 2.29 mm (0.090 in.). (t) To 0.81 mm (0.032 in.). (u) To 0.20 mm (0.008 in.). (v) Shot peening is recommended for wire diameter greater than 1.57 mm (0.062 in.) and smaller where obtainable. (w) To 3.94 mm (0.155 in.). (x) To 7.77 mm (0.306 in.). (y) To 4.76 mm (0.187 in.). (z) To 2.34 mm (0.092 in.). (aa) To 3.76 mm (0.148 in.). (bb) To 9.19 mm (0.362 in.). (cc) To 1.57 mm (0.062 in.). (dd) To 4.50 mm (0.177 in.). (ee) To 0.74 mm (0.029 in.). (ff) To 0.36 mm (0.014 in.).

Fatigue strength is another important mechanical property of steel springs. However, this property is affected by many factors, and because of this complexity, fatigue is discussed in a separate section of this article. Flat Springs. Figure 3 illustrates the different working stresses allowable in flat and leaf springs of 1095 steel that are to be loaded in each of three different ways: statically, variably, and dynamically. (These three types of loading are dealt with separately in the selection table for cold-wound springs, Table 2 ). The stresses given in Fig. 3 are the maximum stresses expected in service. These data apply equally well to 1074 and 1050 steels if the stress values are lowered 10 and 20%, respectively. Except for motor or power springs and a few springs involving only moderate forming, most flat springs, because of complex forming requirements, are formed soft, then hardened and tempered. Fig. 3 Maximum working stress for bending flat and leaf springs made of 1095 steel

For the optimum combination of properties, hypereutectoid spring steel in coil form should be held at hardening temperature for the minimum period of time. The presence of undissolved carbides indicates proper heat treatment. The extent of decarburization can be determined by microscopic examination of transverse sections or by microhardness surveys using Vickers


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or Knoop indentors with light loads (usually 100 g). Power (clock) springs made from pretempered stock have a longer service life if controlled heat treatment can produce a fine, tempered martensitic structure with uniform distribution of excess carbide. If carbides are absent and the tempered martensitic structure is relatively coarse grained, the springs will have a smaller maximum free diameter after having been tightly wound in the barrel or retainer for a long time. A recent development for lower-stressed flat springs is a hardened, 0.04 to 0.22% plain carbon strip steel, which is blanked, formed, and used with only a low-temperature stress-relief treatment. Thickness tolerances, however, are not as close as for spring steel. This material is available in tensile strengths of 900 to 1520 MPa (130 to 220 ksi).

Characteristics of Spring Steel Grade General Spring Quality Wire. The three types of wire used in the greatest number of applications of cold-formed springs are: • Hard-drawn spring wire • Oil-tempered wire • Music wire Hard-Drawn Spring Wire (ASTM A 227). Among the grades of steel wire used for cold-formed springs (Table 1 ), hard-drawn spring wire is the least costly. Its surface quality is comparatively low with regard to such imperfections as hairline seams. This wire is used in applications involving low stresses or static conditions. Oil-tempered wire (ASTM A 229) is a general-purpose wire, although it is more susceptible to the embrittling effects of plating than hard-drawn spring wire. Its spring properties are obtained by heat treatment. Oil-tempered wire is slightly more expensive than hard-drawn wire; it is significantly superior in surface smoothness, but not necessarily in seam depth. Most cold-wound automotive springs are made of oil-tempered wire, although a small percentage are made of music wire and hard-drawn spring wire. Music wire (ASTM A 228) is the carbon steel wire used for small springs. It is the least subject to hydrogen embrittlement by electroplating (see the section "Plating of Springs" in this article) and is comparable to valve-spring wire in surface quality. Chromium-silicon and chromium-vanadium steel spring wire and strip are suitable for moderately elevated temperature service. The chromium-silicon steel spring wire, which has better relaxation resistance than the chromium-vanadium alloy, can be used at temperatures as high as 230 °C (45 °F). The cold-drawn spring wires of the chromium-vanadium and chromium-silicon alloys (A STM A 231 and A 401, respectively) are heat treated before fabrication, while cold-rolled chromium-vanadium (AMS 6455) and chromium-silicon (AISI 9254) strip steels (and generally carbon strip steel as well) are heat treated after rolling and spring fabrication. The chromium-vanadium and chromium-silicon steel spring wires (ASTM A 231 and A 401) can be in either the annealed or oil-tempered condition before spring fabrication. Annealing can be performed before and after drawing, while oil tempering is performed after cold drawing. High-tensile hard-drawn wire fills the gap where high strength is needed but where the quality of music wire is not required. Valve-Spring Quality (VSQ) Wire. All valve-spring wires have the highest surface quality attainable in commercial production, and most manufacturers require that the wire conform to aircraft quality as defined in the AISI Steel Products Manual. Most VSQ wire producers remove the surface of the wire rod before drawing to final size. This practice improves the surface quality and eliminates decarburization. Carbon steel spring wire is the least costly of the VSQ wires. The requirements for carbon VSQ wire in an oil-tempered condition are specified in ASTM A 230. Chromium-vanadium steel wire of valve-spring quality (ASTM A 232) is superior to the same quality of carbon steel wire (ASTM A 230) for service at 120 °C (250 °F) and above. A modified chromium-vanadium steel of valve-spring quality is also specified in ASTM A 878. This modified chromium-vanadium wire has a smaller range of preferred diameters than ASTM A 232 and a lower minimum and maximum tensile strength than ASTM A 232 for a given wire diameter. Chromium-silicon steel VSQ wire (ASTM A 877) can be used at temperatures as high as 230 °C (450 °F). The elevated-temperature behavior of this and other steel spring wires is discussed in the section "Effect of Temperature" in this article. Annealed Spring Wire. Carbon steel wire of valve-spring quality, as well as chromium-vanadium and chromium-silicon steel wire of both spring and valve-spring quality, can be supplied in the annealed condition. This will permit severe forming of springs with a low spring index (ratio of mean coil diameter to wire diameter) and will also permit sharper bends in end hooks. (Although a sharp bend is never desired in any spring, it is sometimes unavoidable.) Springs made from annealed wire can be quenched and tempered to spring hardness after they have been formed. However, without careful control of processing, such springs will have greater variations in dimensions and hardness. This method of making springs is usually used only for springs with special requirements, such as severe forming, or for small quantities, because springs made by this method may have less uniform properties than those of springs made from pretempered wire and are higher in cost. The amount of cost increase depends largely on design and required tolerances, but the cost of heat treating (which often involves fixturing expense) and handling can increase total cost by more than 100%. Stainless Steel Spring Wire. Cold-drawn type 302 stainless steel spring wire (specified in ASTM A 313) is high in heat


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resistance and has good corrosion resistance. The surface quality of type 302 stainless steel spring wire occasionally varies, seriously affecting fatigue resistance. Type 316 stainless is superior in corrosion resistance to type 302, particularly against pitting in salt water, but is more costly and is not considered a standard spring wire. Type 302 is readily available and has excellent spring properties in the full-hard or spring-temper condition. It is more expensive than any of the carbon steel wires for designs requiring a diameter larger than about 0.30 mm (0.012 in.) but is less expensive than music wire for sizes under about 0.30 mm (0.012 in.). In many applications, type 302 stainless can be substituted for music wire with only slight design changes to compensate for the decrease in modulus of rigidity. For example, a design for a helical compression spring was based on the use of 0.25 mm (0.010 in.) diam music wire. The springs were cadmium plated to resist corrosion, but they tangled badly in the plating operation because of their proportions. A redesign substituted type 302 stainless steel wire of the same diameter for the music wire. Fewer coils were required because of the lower modulus of rigidity, and the springs did not require plating for corrosion resistance. The basic cost of this small-diameter stainless wire was, at the time, 20% less than the cost of the music wire. Elimination of plating and reduction of handling resulted in a total savings of 25%.

Wire Quality Specification requirements for the spring materials listed in Table 1 include twist, coiling, fracture, or reduction-in-area tests, in addition to dimensional limits and minimum tensile strength. Such tests ensure that the wire has the expected ductility and has not been overdrawn (which could produce internal splits or voids). In dynamic applications, in which fatigue strength is an important factor, the performance differences of spring materials depends on surface quality when materials are of similar composition and tensile strength. Because the initiation and growth of fatigue cracks is strongly affected by surface quality, seams and surface decarburization are important factors in dynamic applications of spring quality wire and especially valve-spring quality wire. Freedom from surface imperfections is of paramount importance in some applications of highly stressed springs for shock and fatigue loading, especially where replacement of a broken spring would be difficult and much more costly than the spring itself or where spring failure could cause extensive damage to other components. Seams are evaluated visually, often after etching with hot 50% muriatic acid. The depth of metal removed can vary from 0.006 mm (1=4mil) to 1% of wire diameter. Examination of small-diameter etched wire requires a stereoscopic microscope, preferably of variable power so that the sizes of seams can be observed in relation to the diameter of the wire. The least expensive wires can have seams that are quite pronounced. Hard-drawn and oil-tempered wires occasionally have seams as deep as 3.5% of wire diameter, but usually not deeper than 0.25 mm (0.010 in.). On the other hand, wires of the highest quality (music wire, valve-spring) have only small surface imperfections, generally not deeper than 1=2% of wire diameter. Some grades can be obtained at moderate cost, with seam depth restricted to 1% of wire diameter. Decarburization. There is no general numerical limit on decarburization, and phrases such as "held to a minimum consistent with commercial quality" are very elastic. It is usual for seams present during hot rolling to be partly decarburized to the full depth of the seam or slightly deeper. For valve-spring quality, however, decarburization limits are more severe. Some manufacturers permit loss of surface carbon only it if does not drop below 0.40% for the first 0.025 mm (0.001 in.) and, within the succeeding 0.013 mm (0.0005 in.), becomes equal to the carbon content of the steel. As noted previously, most manufacturers of VSQ wire eliminate surface decarburization by removing the surface of the wire rod prior to drawing to final size. General decarburization can be detrimental to the ability to maintain load. For hot-wound springs made directly from hot-rolled bars, it is common practice to specify a torsional modulus of 72 GPa (10 × 106 psi) instead of the 80 GPa (11.5 × 106 psi) used for small spring wires. In part, this compensates for the low strength of the surface layer. Total loss of carbon from the surface during a hear-treating process is infrequent in modern wire mill products. Partial decarburization of spring wire is often blamed for spring failures, but quench cracks and coiling-tool marks are more frequently the actual causes. In wires of valve-spring or aircraft quality, a decarburized ferritic ring around the wire circumference is a basis for rejection. The net effects of seams and decarburization are described in the section "Fatigue " in this article. Magnetic Particle and Eddy Current Testing. Inspection for seams and other imperfections in finished springs is generally carried out by magnetic particle inspection. In its various forms, this inspection method has proved to be the most practical nondestructive method for the inspection of springs that may affect human safety or for other reasons must not fail as a result of surface imperfections. For compression and extension springs, the inspection is always concentrated selectively on the inside of the coil, which is more highly stressed than the outside and is the most frequent location of start of failure. Valve-spring wire is often 100% eddy current tested for seams and point defects by the wire mill. The defects are painted to ensure that they are not fabricated into finished springs.

Residual Stresses Residual stresses can increase or decrease the strength of a spring material, depending on their direction. For example, residual stresses induced by bending strengthen wire for deflection in the same direction while weakening it for deflection in the opposite direction. In practice, residual stresses are either removed by stress relieving or induced to the proper direction by cold setting and shot peening. Compression springs, torsion springs, flat springs, and retaining rings can be stress relieved and cold set. The treatment used


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depends on the design and application requirements of the individual spring. Many compression springs are preset for use at higher stress. They are then known as springs with set removed. Compression springs can be made to close solid, without permanent set and without presetting, if the shear stress is less than a specified proportion of the tensile strength of the wire in the fully compressed spring (about 45% for music wire) and if the springs are properly stress relieved. The maximum shear stress in a fully compressed preset spring is about 33% higher, or approximately 60% of the tensile strength of the wire. Therefore, presetting to a maximum stress will permit the use of up to 45% less steel than is otherwise required, a savings greater than the cost of presetting for wire larger than about 3 mm (1=8in.) in diameter. Also, the smaller, equally strong spring requires less space. When the calculated uncorrected stress at solid height is greater than about 60% of the tensile strength (or, for cold-set springs, greater than the proportional limit stress), the spring can be neither cold set nor compressed to its solid height without taking a permanent set. Several types of springs are in this category, where the maximum permissible deflection must be calculated and positive stops provided to avoid permanent set in service. Compression springs, cold wound from pretempered or hard-drawn high-carbon spring wire, should always be stress relieved to remove residual stresses produced in coiling. Extension springs are usually given a stress-relieving treatment to relieve stresses induced in forming hooks or other end configurations, but such treatment should allow retention of stresses induced for initial tension. The treatment of wire retaining rings depends on whether the loading tends to increase or decrease the relaxed diameter of the spring. Most rings contain residual stresses in tension on the inside surface. For best performance, rings that are reduced in size in the application should not be stress relieved, while expanded rings should be. This consideration applies equally to torsion springs. It is common practice to give these springs a low-temperature heat treatment to provide dimensional stability. Stress relieving affects the tensile strength and elastic limit, particularly for springs made from music wire and hard-drawn spring wire. The properties of both types of wire are increased by heating in the range of 230 to 260 °C (450 to 500 °F). Oil-tempered spring wire, except for the chromium-silicon grade, shows little change in either tensile strength or elastic limit after stress relieving below 315 °C (600 °F). Both properties then drop because of temper softening. Wire of chromium-silicon steel temper softens only above about 425 °C (800 °F). The properties of spring steels are usually not improved by stress relieving for more than 30 min at temperature, except for age-hardenable alloys such as 631 (17-7 PH) stainless steel, which requires about 1 h to reach maximum strength. Typical stress-relief temperatures for steel spring wire are given in Table 3 . Table 3 Typical stress-relieving temperatures for steel spring wire Applicable only for stress relieving after coiling and not valid for stress relieving after shot peening Temperature(a) °C

°F

Music wire

230−260

450−500

Hard-drawn spring wire

230−290

450−550

Oil-tempered spring wire

230−400

450−750(b)

Valve spring wire

315−400

600−750

Cr-V spring wire

315−400

600−750

Cr-Si spring wire

425−455

800−850

Type 302 stainless

425−480

800−900

Steel

Type 631 stainless 480 ± 6(c) 900 ± 10(c) (a) Based on 30 min at temperature. (b) Temperature is not critical and can be varied over the range to accommodate problems of distortion, growth, and variation in wire size. (c) Based on 1 h at temperature

When springs are to be used at elevated temperatures, the stress-relieving temperatures should be near the upper limit of the range to minimize relaxation in service. Otherwise, lower temperature are better.

Plating of Springs Steel springs are often electroplated with zinc or cadmium to protect them from corrosion and abrasion. In general, zinc has been found to give the best protection in atmospheric environments, but cadmium is better in marine and similar environments involving strong electrolytes. Electroplating increases the hazards of stress raisers and residual tensile stresses because the hydrogen released at the surface during acid or cathodic electrocleaning or during plating can cause a time-dependent brittleness, which can act as though added tensile stress had been applied and can result in sudden fracture after minutes, hours, or hundreds of hours. Unrelieved tensile stresses can result in fracture during plating. Such stresses occur most severely at the inside of small-radius bends. Parts with such bends should always be stress relieved before plating. However, because even large-index springs have been found to be cracked, general stress relief is always good practice. Preparation for plating is also very important because hydrogen will evolve from any inorganic or organic material on the metal until the material is thoroughly covered. Such contaminants may be scarcely noticeable before plating except by their


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somewhat dark appearance. Thorough sandblasting or tumbling may be required to remove such layers. Hydrogen Relief Treatment. If stress relieving has been attended to, and the springs are truly clean before plating, then the usual baking treatment of around 200 °C (400 °F) for 4 h should lessen the small amount of hydrogen absorbed and redistribute it to give blister-free springs, which will not fail. Mechanical Plating. Another technique that solves the hydrogen problem is mechanical plating, which involves cold welding particles of zinc or other soft-metal powder to an immersion copper flash plate on the spring. While some hydrogen may be absorbed during acid dipping before plating, it does not result in a time-dependent embrittlement because the plated layer is inherently porous, even though it has a shiny appearance. The hydrogen easily diffuses through the pores within 24 h, leaving the steel ductile.

Fatigue For those springs that are dynamically loaded, it is common practice to obtain basic mechanical data from S-N fatigue curves. A typical S-N diagram is shown in Fig. 4 (a). For each cycle of fatigue testing, the minimum stress is zero and the maximum stress is represented by a point on the chart. An alternative method of presenting data on the fatigue life of springs is shown in Fig. 4 (b). Fig. 4 Fatigue lives of compression coil springs made from various steels. (a) S-N diagram for springs made of minimum-quality music wire 0.56 mm (0.022 in.) in diameter. Spring diameter was 5.21 mm (0.205 in.); D/d was 8.32. Minimum stress was zero. Stresses corrected by Wahl factor (see the section "Wahl Correction" in this article). (b) Life of springs used in a hydraulic transmission. They were made of oil-tempered wire (ASTM A 229) and music wire (ASTM A 228). Wire diameter was 4.75 mm (0.187 in.), outside diameter of spring was 44.45 mm (1.750 in.), with 15 active coils in each spring. The springs were fatigue tested in a fixture at a stress of 605 MPa (88 ksi), corrected by the Wahl factor.

Stress Range. In most spring applications, the load varies between initial and final positive values. For example, an automotive valve spring is compressed initially during assembly, and during operation it is further compressed cyclically each time the valve opens. The shear-stress range (that is, the difference between the maximum and minimum of the stress cycle to which a helical steel spring may be subjected without fatigue failure) decreases gradually as the mean stress of the loading cycle increases. The allowable maximum stress increase up to the point where permanent set occurs. At this point, the maximum stress is limited by the occurrence of excessive set. Figure 5 shows a fatigue diagram for music wire springs of various wire diameters and indexes. This is a modified Goodman diagram and shows the result of many fatigue-limit tests on a single chart. In Fig. 5 , the 45° line OM represents the


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minimum-stress of the cycle, while the plotted points represent the fatigue limits for the respective minimum stresses used. The vertical distances between these points and the minimum-stress reference line represent the stress ranges for the music wire springs. Fig. 5 Fatigue limits for compression coil springs made of music wire. Data are average fatigue limits from S-N curves for 185 unpeened springs of various wire diameters run to 10 million cycles of stress. All stresses were corrected for curvature using the Wahl correction factor. The springs were automatically coiled, with one turn squared on each end, then baked at 260 °C (500 °F) for 1 h, after which the ends were ground perpendicular to the spring axis. The test load was applied statically to each spring and a check made for set three times before fatigue testing. The springs were all tested in groups of six on the same fatigue testing machine at ten cycles per second. After testing, the unbroken springs were again checked for set and recorded. Number 4 springs, tested at 1070 MPa (155 ksi) maximum stress, had undergone about 21=2% set after 10 million stress cycles, but the stresses were not recalculated to take this into account. None of the other springs showed appreciable set. The tensile strengths of the wires were according to ASTM A 228.

Wire diameter

Spring outside diameter

Spring No.

mm

in.

mm

in.

Spring index

1

0.81

0.032

9.52

0.375

10.7

2

0.81

0.032

6.35

0.250

6.8

3

1.22

0.048

15.88

0.625

12.0

4

2.59

0.102

22.22

0.875

7.6

5

3.07

0.121

22.22

0.875

6.2

450

0.177

22.22

0.875

4.9

6

Free length mm

in.

Total turns

Active turns

Total tested

22.10

0.87

6.0

4.2

16

26.97

1.062

7.0

5.2

28

44.45

1.75

7.0

5.2

38

60.20

2.37

7.0

5.2

43

57.15

2.25

7.5

5.7

35

57.15

2.25

7.5

5.7

25

In fatigue testing, some scatter may be expected. The width of the band in Fig. 5 may be attributed partly to the normal changes in tensile strength with changes in wire diameter. There appears to be a trend toward higher fatigue limits for the smaller wire sizes. Line UT is usually drawn so as to intersect line OM at the average ultimate shear strength of the various sizes of wire. Modified Goodman diagrams for helical springs made of several steels are shown in Fig. 6 (music wire and 302 stainless steel wire) and Fig. 7 (hard-drawn steel spring wire, oil-tempered wire, and chromium-silicon steel wire). In all instances, the plotted stress values were corrected by the Wahl factor (see the section "Wahl Correction" in this article). The data were obtained from various sources, including controlled laboratory fatigue tests, spot tests on production lots of springs, and correlation between


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rotating-beam fatigue tests on wire and uni-directional-stress fatigue tests on compression and extension helical springs. Fig. 6 Modified Goodman diagrams for steel helical springs made from music wire (a and b) and 302 stainless steel wire (c and d). The graphs on the left (a and c) plot maximum allowable stresses for 10 million cycles for a similar group of wire diameters. All stresses were corrected by the Wahl factor. See text for discussion.

Fig. 7 Modified Goodman diagrams for steel helical springs made from chromium-silicon steel (a and b), oil-tempered wire (c and d), and hard-drawn spring wire (e and f). The graphs on the left (a, c, and e) plot maximum allowable stress for 10 million cycles for 3.18 mm (0.125 in.) diam wires and various other size wires. All stresses were corrected by the Wahl factor. See text for discussion.


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The graphs on the left side of Fig. 6 and 7 plot the allowable stresses at 10 million cycles taken from S-N curves for various wire size diameters. The graphs on the right side of Fig. 6 and 7 show the allowable stresses at 10,000, 10,000, and 10 million cycles for two different wire diameters. In Fig. 6 and 7 , the stress range is the vertical distance between the 45° line and the


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lines for the several wire sizes. The allowable maximum stress increases to a point of permanent set, indicated by the horizontal sections of the lines on the diagrams. For equal wire sizes, these diagrams show that music wire (Fig. 6 a and b) is the most fatigue-resistant wire, with fatigue limits 50% greater than those of the least resistant wires (hard-drawn wires, Fig. 7 e and f). This difference is largely maintained under high-stress, short-life conditions. Figures 6 and 7 represent normal quality for each grade. Because of variations in production conditions, however, quality is not constant. Figure 8 shows the range of fatigue life for five lots of music wire, all 0.50 mm (0.020 in.) in diameter. Wire was tested on a rotating-beam machine at a maximum stress of 1170 MPa (170 ksi) and a mean stress of zero. Results were correlated with fatigue tests on torsion springs as follows. A minimum fatigue life of 50,000 cycles was required of each spring; a minimum life of 20,000 cycles for the wire in the rotating-beam machine at 1170 MPa (170 ksi) gave satisfactory correlation with the 50,000 cycle service life of springs made from the wire. Lot 5 in Fig. 8 was rejected because it failed to meet the fatigue requirement. Subsequent fatigue tests on a pilot lot of springs made from lot 5 wire confirmed the inability of these springs to meet the fatigue requirement of 50,000 cycles. Fig. 8 Fatigue-life distribution of 0.50 mm (0.020 in.) diam music wire. Tested in a rotating-beam machine at a maximum stress of 1170 MPa (170 ksi) and a mean stress of zero

Shot peening of springs improves fatigue strength by prestressing the surface in compression. It is usually applied to wire 1.6 mm (1=16in.) or more in diameter. The type of shout used is important; better results are obtained with carefully graded shot having no broken or angular particles. Shot size may be optimum at 10 to 20% of the wire diameter. However, for larger wire, it has been found that excessive roughening during peening with coarse shot lessens the benefits of peening, apparently by causing minute fissures. Also, peening thin material too deeply leaves little material in residual tension in the core; this negates the beneficial effect of peening, which requires internal tensile stress to balance the surface compression. Shot peening is effective in largely overcoming the stress-raising effects of shallow pits and seams. Proper peening intensity is an important factor, but more important is the need for both the inside and outside surfaces of helical springs to be thoroughly covered. An Almen test strip necessarily receives the same exposure as the outside of the spring, but to reach the inside, the shot must pass between the coils and is thus greatly restricted. As a result, for springs with closely spaced coils, a coverage of 400% on the outside may be required to achieve 90% coverage on the inside. Cold-wound steel springs are normally stress relieved after peening to restore the yield point. A temperature of 230 °C (450 °F) is common because higher temperatures degrade or eliminate the improvement in fatigue strength. The extent of improvement in fatigue strength to be gained by shot peening, according to one prominent manufacturer of cold-wound springs, is shown in Fig. 9 . The bending stresses apply to flat springs, power springs, and torsion springs; the torsional stresses apply to compression and extension springs. Fig. 9 Fatigue curves for peened and unpeened steel spring wires


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Effect of Temperature The effect of elevated temperatures on the mechanical properties and performance of fabricated springs is shown in Fig. 10 , 11 , 12 , 13 , and 14 . The effect is reported as amount of load loss (relaxation), which is a function of chemical composition, maximum stress, and spring processing. Fig. 10 Relaxation curves for steel helical springs of music wire (ASTM A 228), chromium-silicon spring wire (ASTM A 401), oil-tempered spring wire (ASTM A 229), chromium-vanadium spring wire (ASTM A 231), and hard-drawn spring wire (ASTM A 227) at (a) 90 °C (200 °F) and (b) 150 °C (300 °F). Relaxation curves for the low-alloy steel spring wires (ASTM A 231 and A 401) are also plotted at (c) 200 °C (400 °F) and (d) 260 °C (500 °F). All curves represent relaxation after exposure for 72 h at indicated temperatures.


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Fig. 11 Relaxation curves for steel helical springs made of (a) 302 stainless steel and (b) 631 stainless steel. The curves represent relaxation after exposure for 72 h at the indicated temperatures.

Fig. 12 Effect of time and temperature on the relaxation of ten-turn helical springs made from (a) music wire per ASTM A 228 and (b) 420 and 431 stainless steel wire. Wire diameter, 2.69 mm (0.106 in.); spring diameter, 25.4 mm (1.00 in.); free length, 76.2 mm (3.00 in.). Stresses were corrected by the Wahl factor.


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Fig. 13 Relaxation of carbon steel (1070, 1095) and SAE 52100 alloy steel circular flat springs (piston rings) at elevated temperatures. Spring hardness was 35 HRC. Springs were exposed to the indicated temperatures for 3 to 4 h.

Fig. 14 Load-loss curves at various temperatures for helical compression steel springs made of music wire (ASTM A 228), VSQ carbon steel spring wire (ASTM A 230), VSQ chromium-vanadium steel spring wire (ASTM A 232), type 302 stainless steel spring wire (ASTM A 313), VSQ chromium-silicon steel spring wire (AISI 9254), and T1 high-speed tool steel. Results are based on tests of thousands of springs. All springs were made of pretempered wire and were stress relieved after coiling; none were shot peened, and all curves are for exposure of at least 72 h at the indicated temperatures.


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Figures 10 and 11 show some results from testing helical compression springs to determine the maximum relaxation at a given static working stress, temperature, and time. A specific spring height was determined for a given corrected stress. The spring was clamped at this height and placed in a convection over for 72 h. It was removed and cooled, and the new free height was measured. Load loss was determined and amount of relaxation calculated. For example, a music wire spring designed at a corrected stress of 690 MPa (100 ksi) will relax a maximum of 3.8% when held at 90 °C (200 °F) for 72 h. From the results shown in Fig. 10 (a) and 10 (b), springs made from music wire (ASTM A 228) are equal in performance to those made from oil-tempered wire (ASTM A 229) at 90 °C (200 °F) but are inferior at 150 °C (300 °F). The springs made from low-alloy steel wire such as chromium-silicon (ASTM A 401) or chromium-vanadium (ASTM A 231) provide improved relaxation resistance at elevated temperatures (Fig. 10 ), while the stainless steel springs (Fig. 11 ) exhibit further improvements in relaxation resistance. Percentage of load loss increases with shear stresses. The effect of time and temperature on relaxation of springs is shown in Fig. 12 . Rate of relaxation is greatest during the first 50 to 75 h. For longer periods of time, the rate is lower, decreasing as the logarithm of time when no structural change or softening occurs. A significant effect of temperature on the relaxation of piston rings is indicated in Fig. 13 . The rings were confined in test cylinders to maintain the outside ring diameter and were exposed to test temperature for 3 to 4 h. The load required to deflect the ring to working diameter was measured before and after each test to calculate the amount of relaxation. Tests on thousands of springs under various loads at elevated temperatures are summarized in Fig. 14 . Plain carbon spring steels of valve-spring quality are reliable at stresses up to 550 MPa (80 ksi) (corrected) and temperatures no higher than 175 °C (350 °F), in wire sizes no greater than 9.5 mm (3=8in.). Slightly more severe applications may be successful if springs are preset at the operating temperature with loads greater than those of the application. Plain carbon spring steels of valve-spring quality should not be used above 200 °C (400 °F). To achieve an acceptable balance between relaxation and stress level, it is recommended that plain carbon steels be used at 120 °C (250 °F) or below (Table 1 ). Except for high-speed steel, these tests revealed no advantage in springs heat treated after coiling compared with those at the same hardness made of pretempered wire and properly stress relieved. Deflection of a spring under load is inversely proportional to the modulus of rigidity, G, of the material. Variation with temperature is shown in Fig. 15 . Fig. 15 Effect of temperature on modulus of rigidity of spring steels


Page 519


Steel Springs

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A uniform deflection under load over a range of temperatures sometimes must be maintained. The instrument spring shown in Fig. 16 required a constant modulus up to 90 °C (200 °F) and, when made of music wire, drifted 5% in service. It was replaced with a satisfactory spring made of a nickel alloy with constant modulus (Fe-42Ni-5.4Cr-2.40Ti-0.60Al-0.45Mn-0.55Si-0.06C). Fig. 16 Helical instrument spring that required a constant modulus of rigidity up to 90 °C (200 °F)

In another example, a spring was originally fabricated from oil-tempered wire (ASTM A 229) and performed satisfactorily when tested a room temperature. However, in service it was immersed in oil that attained a temperature slightly above 90 °C (200 °F), which was high enough to cause excessive relaxation over a period of 2 to 3 h. Chromium-silicon steel spring wire (ASTM A 401) was substituted for the oil-tempered wire, and at the identical operating temperature service was satisfactory. Design data for these springs are given in Table 4 . Table 4 Comparison of two spring materials for elevated-temperature service Condition at elastic limit Load Grade of wire(a) Oil-tempered (A229)

Total deflection

kg

lb

mm

in.

122.7

270.5

19.3

0.76

Cr-Si steel (A401) 139.7 308.0 21.8 0.86 (a) Design data (the same for both types of wire): mean diameter, 15.88 mm (0.625 in.); inside diameter, 12.12 mm (0.477 in.); wire diameter; 3.76 mm (0.148 in.); spring index, 4.22; total number of coils, 10; number of active coils, 8; spring rate, 6.36 kg/mm (356 lb/in.); working load, 5.17 kg (114 lb); deflection at working load, 7.92 mm (0.312 in.); stress at working load, 383 MPa (55.6 ksi); solid height, 37.6 mm (1.48 in.) max; free height, ~49.0 mm (1.93 in.); set removed, plain finish, variable type of load, environment of about 90 °C (200 °F)

Hot-Wound Springs Although some hot-wound springs are made of steels that are also used for cold-wound springs, hot-wound springs are usually much larger, which results in significant metallurgical differences. Hardenability Requirements. Steels for hot-wound springs are selected mainly on the basis of hardenability. Carbon steels with about 0.70 to 1.00% C (1070 to 1095) are suitable and widely used for statically and dynamically loaded springs in the smaller sizes. Carbon steels are also used for larger springs in the lower stress range, where some hardenability can be sacrificed safely. In most cases, alloy steels are usually required for the larger sizes because of the need for hardenability. Most specifications for hot-wound alloy steel springs require 0.50 to 0.65% C and a minimum hardness of 50 HRC at the center after oil quenching from about 815 °C (1500 °F) and before tempering. (The austenitizing temperature will vary, depending on the specific steel.) Springs subjected to lower stress ranges may not require this high hardenability and can therefore be made of the lower-priced


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01 Sep 2005

carbon steels. Typical minimum hardness and hardenability values are given in Table 5 for hot-wound springs used under specific conditions of stress. Here it can be seen that lower as-quenched surface hardness and lower hardenability are permitted in the lower stress ranges for both statically and variably loaded springs. These requirements gradually increase as the stress increases. After tempering, a hardness of 50 HRC is often specified for a static stress of 1240 MPa (180 ksi) or for springs dynamically loaded at 825 MPa (120 ksi). Table 5 Typical minimum hardness and hardenability for steel used for hot-wound helical springs As-oil-quenched, prior to tempering. Normal hardness, as-tempered, 44−49 HRC at surface Corrected maximum solid stress MPa

Hardness, HRC

ksi

At surface

At center

690

100

45

35

825

120

50

45

Static load

965

140

60

50

1100

160

60

50

1240

180

60

50

Variable load, designed for a minimum life of 50,000 cycles (set removed, 2.5% probability of failure, mean stress 515 MPa, or 75 ksi) 690

100

45

35

825

120

50

45

965

140

60

50

1100

160

60

50

1240

180

60

50

Dynamic load, designed for a minimum life of 2 million cycles (set removed, shot peened, 2.5% probability of failure, mean stress 515 MPa, or 75 ksi) 690

100

45

35

825

120

60

50

Recommended steels for hot-wound helical springs are given in Table 6 , covering variations in stress range, type of loading, and wire size. Hardenability requirements increase as required strength and/or wire diameter increases. Table 6 Recommended steels for hot-wound helical springs Where more than one steel is recommended for a specific set of conditions, they are arranged in the order of increasing hardenability. The first steel listed applies to the lower end of the designated wire diameter range and the last to the upper end of the range. Corrected maximum solid stress Diameter of spring wire (hot rolled) MPa

ksi

9.5−25.4 mm (3=8−1 in.)

25.4−50.8 mm (1−2 in.)

50.8−76.2 mm (2−3 in.)

1095

1095

Static load (set removed, static stress up to 80% of maximum solid stress) 825

120

1070 1095

965

140

1070 1095

1100

160

1240

180

51B60H 4161

4161

5150H 5160H 50B60H

51B60H 4161

4161

5150H 5160H 50B60H

51B60H 4161

4161

Variable load, designed for minimum life of 50,000 cycles (set removed, 2.5% probability of failure, operating stress range not over 50% of solid stress) 690

100

1095

825

120

1095

965

140


5150H 5160H

1095

1095

51B60H 4161

4161

51B60H 4161

4161

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Steel Springs

01 Sep 2005

50B60H 1100

160

5150H 5160H 50B60H

51B60H 4161

4161

1240

180

5150H 5160H 50B60H

51B60H 4161

4161

Dynamic load, designed for minimum life of 2 million cycles (set removed, shot peened, 2.5% probability of failure, mean stress 515 MPa, or 75 ksi, operating stress range not over 50% of solid stress) 690

100

1095

825

120

1095

1095 51B60H 4161

1095 4161

The strengths obtained in bars with different hardenabilities are shown in Fig. 17 . The band for 1095 steel demonstrates the variation in yield strength that results from variations in the thickness of the bar and the severity of quenching and from limited hardenability. Smaller bars, as well as those quenched more rapidly from the austenitizing temperature, follow the top of the band, while larger bars and those quenched less drastically fall in the lower half of the band. For example, a 13 mm (1=2in.) round bar quenched in 5% caustic solution will have maximum properties, and a 50 mm (2 in.) round bar quenched in still oil will have yield strength near the minimum shown. The scatter for alloy steels in Fig. 17 is much narrower because the points only represent steels with sufficient hardenability to have a martensitic structure throughout the bar section, as-quenched. Fig. 17 Relationship between surface hardness and yield strength of steel bars after tempering. Alloy steels were martensitic throughout the section, as-quenched; 1095 bars were 12.7 to 50.8 mm (1=2to 2 in.) in diameter.

The distribution of hardness test results at surface and center, as-quenched and as-tempered, is shown in Fig. 18 for a multiplicity of heats of 1095 steel and five alloy steels commonly used for hot-wound springs. The results of testing specimens 300 mm (12 in.) long correlate with those of testing production springs. For the alloy steels, a minimum of 50% martensite at the center of the quenched section was specified. Table 7 lists steels that meet hardenability requirements for torsion bar springs with section thicknesses from 29.21 to 57.15 mm (1.150 to 2.250 in.). Fig. 18 Hardness distribution for steels for hot-wound helical springs. Alloy steels were oil quenched from 845 °C (1550 °F); 1095 was oil quenched from 885 °C (1625 °F). Data were obtained from hot-rolled, heat-treated laboratory test coupons, 305 mm (12 in.) long. Specimens were sectioned from the center of the coupons after heat treatment. These results on bars correlate with those on production springs. For the alloy steels, a minimum of 50% martensite at the center of the quenched section was specified.


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Table 7 Steels with sufficient hardenability for torsion bar springs Thickness of section

Minimum hardenability required

mm

in.

29.21

1.150

J50 at 8

33.91

1.335

J50 at 9

Steel 8650H, 5152H 8655H, 50B60H

1

36.32

1.430

J50 at 8 =2

50B60H

39.88

1.570

J50 at 10

51B60H

46.23

1.820

J50 at 11

8660H

48.26

1.900

J50 at 14

4150H

57.15

2.250

J50 at 22

9850H


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Steel Springs

01 Sep 2005

Surface Quality. Hot-wound springs fabricated from bars of large diameter will normally show much deeper surface decarburization, in the range from 0.13 to 0.38 mm (0.005 to 0.015 in.), unless special material preparation and processing techniques are used. Although decarburization is not desirable, in some design situations the use of the favorable residual stress pattern due to shallow hardening can reduce the negative effects of decarburization on fatigue performance (Table 8 ). The detrimental effects of decarburization are less noticeable on hot-wound than on cold-wound springs because other weaknesses, such as the surface imperfections and irregularities typical of a hot-rolled surface, create additional focal points for fatigue cracks. Table 8 Properties of shallow-hardened and through-hardened hot-wound (wound from bar 32 mm, or 1 1=4 in., in diameter) helical compression springs Hardness, HRC Distance below surface Heat treatment

1.6 mm (1=16 in.)

Fatigue tests(a) Millions of cycles to failure

3.2 mm (1=8 in.)

6.35 mm (1=4 in.)

Center

High

Low

Average

60

50

45

30

...

...

...

54.5

49

37.5

25

1.28

0.46

0.772

60

60

60

...

...

1045 steel, decarburized 0.38 mm (0.015 in.) As-quenched Tempered

8655H, steel decarburized 0.03 mm (0.001 in.) As-quenched

60

Tempered 42 43 43 42 0.789 0.443 (a) 485 MPa (70 ksi) stress range (corrected), 415 MPa (60 ksi) mean stress (corrected); eight springs tested in each group

... 0.572

Specifications for hot-wound springs and torsion bar springs usually include maximum seam depth. For example, the specifications used by a manufacturer of railway equipment allow seams with a maximum depth of 0.025 mm (0.001 in.) per 1.6 mm (1=16in.) of wire diameter up to 0.41 mm (0.016 in.) for any bar size above 25 mm (1 in.). Another manufacturer allows seam depths of 0.41 mm (0.016 in.) for bars 25 mm (1 in.) in diameter, 0.81 mm (0.032 in.) for bars 25 to 44 mm (1.00 to 1.75 in.) in diameter, and 1.22 mm (0.048 in.) for bars 44 to 63.5 mm (1.75 to 2.50 in.) in diameter. Design Stress. It should be noted that some organizations specify a much more conservative approach to stress than that presented in Table 6 . The Spring Manufacturers Institute has adopted a chart from the Manufacturers Standardization Society of the Valve and Fittings Industry, for essentially static service, that calls for an admittedly uncorrected stress for alloy steel of 760 MPa (110 ksi) for a bar diameter of 13 mm (1=2in.), reducing to 590 MPa (86 ksi) for a bar diameter of 92 mm (35=8in.). Their curve for 1095 steel is roughly 140 MPa (20 ksi) lower. The desirability of conservative design in cyclical service is illustrated in Fig. 19 , in which the minimum stress used was low. Such data on springs hot wound from bars with as-rolled surfaces are limited, and interpretation is therefore difficult. The value of peening, however, is made quite apparent. Surface imperfections can be removed by grinding, and this is normal practice, where the increased cost can be borne, in order to increase reliability at higher stresses. Fig. 19 Effect of peening on the probability of fatigue failure of hot-wound steel springs. Top: 95% probability of failure. Middle: 50% probability of failure. Bottom: 5% probability of failure. Springs were made from 16 to 27 mm ( 5=8to 11=16in.) diam 8650 and 8660 hot-rolled steel and heat treated to between 429 and 444 HB. Springs were shot peened to an average arc height of 0.2 mm (0.008 in.) on the type C Almen strip at 90% visual coverage.


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Steel Springs

01 Sep 2005

With all of the limitations discussed above, there can be both cost and reliability advantages in using 1095 steel at low stresses. For example, railroad freight car springs, which are subject to severe corrosion pitting, are made of 1095 steel and designed for very low stress, and very little difficulty is encountered.

Costs The relative costs of various spring steels in the form of round wire are given in Fig. 20 . Base price may be outweighed by other costs, as indicated in Table 9 , which lists possible extras for two alloy spring steels of about the same base price. Also, base price and other costs will vary somewhat with time. Fig. 20 Relative cost of spring steel wire

Table 9 Pricing of automotive coiled spring steel Cost per 45 kg (100 lb)(a)

Item Alloy steel 6150 Base price

Alloy steel


$25.60

Page 525


Steel Springs

01 Sep 2005

Grade

6150

6.75

Quality

Electric furnace

2.30

Restrictions

Restricted carbon and manganese

5.50

Size

13 mm (1=2 in.) diam

3.50

Straightness

Special straightness

3.00

Treatment

Hot rolled, precision ground

8.95

Cleaning and coating

Pickled and oiled

4.10

Preparation

None

Testing

...

Restricted hardenability

3.85

Cutting

1.5−1.4 m (5−8 ft) abrasive

1.00

Length

4.6 m (15 ft) dead length

0.50

Quantity

127(i)

46 429 min min(j)

5150H 5155H 50B44H 5147H 9260H 81B45H 8650H 86B45H 6150H

5160H 50B50H 9262H 4147H 8655H

50B60H 51B60H 8660H

4150H

>185

4145H 9840H

86B45H 4337H

4147H 4340H

4150H

4340H

9805H E4340H

9850H

Water quenched and tempered(k) 620−860( 90−125 a)

23−30( 241−28 b) 5

5130H 5132H 4130H 8630H

5135H

4135H

94B30H

860−1030 125−150 (c)

30−36( 285−34 d) 1

1330H 5135H

1335H 4135H(l) 1340H(m ) 8640H(l) 8740H(l) 8637H(m ) 3140H(l)

50B40H 8642H 94B30H

4137H 4140H

94B40H

1030−117 150−170 0(e)

36−41(f 331−37 ) 5

1330H 1335H 5130H 5132H 5135H 4130H 8630H

4042H 4047H

50B40H(l 8640H(m ) ) 4137H(l) 8740H(m ) 8642H(l) 8745H(l)

50B44H 5147H 4140H 8645H 8742H

94B40H

81B45H 4142H 4337H

1170−127 170−185 5(g)

41−46( 375−42 h) 9

5140H 4037H 4042H 4137H 8637H

1340H 50B46H 3140H

5145H 50B44H(l 4140H(m ) ) 50B40H 8640H 5147H(l) 8645H(m ) 8642H 81B45H(l ) 8742H(m 8740H ) 94B40H(l )

4142H

81B45H 4337H

4145H 4147H 86B45H 9840H 4340H E4340H

1340H 50B46H 5140H 4135H 8637H 94B30H 3140H

81B45H( 4147H 5147H 50B44H 5046H m) 4145H 50B46H 8645H 5145H 4047H 86B45H 4142H 8642H (a) Tensile strength, 790 to 940 MPa (115 to 138 ksi). (b) As-quenched hardness, 42 HRC, or 388 HB. (c) Tensile strength, 940 to 1100 MPa (136 to 160 ksi). (d) As-quenched hardness, 44 HRC, or 415 HB. (e) Tensile strength, 1100 to 1300 MPa (160 to 188 ksi). (f) As-quenched hardness, 48 HRC, or 461 HB. (g) Tensile strength, 1300 to 1530 MPa (188 to 222 ksi). (h) As-quenched hardness, 51 HRC, or 495 HB. (i) Tensile strength, over 1530 MPa (222 ksi). (j) As-quenched hardness, 55 HRC, or 555 HB. (k) Through steels with 0.47% C nominal. (l) May be substituted for steels listed under the 50 to 63 mm (2 to 21=2in.) column at same strength level or less. (m) Not recommended for applications requiring 80% martensite at midradius in sections 38 to 50 mm (11=2to 2 in.) in diameter because of insufficient hardenability. Source: Ref 3 >1275(i)

>185

46 429 min min(j)

Table 4 Alloy steel selection guide for moderately stressed parts Unless otherwise indicated in the footnotes, any steel in this table may be considered for a lower strength level or a smaller section, or both. Steels to give 50% martensite, minimum, for indicated location in a round


Page 753


Hardenable Carbon and Low-Alloy Steels

As-tempered hardness

Required yield strength

MPa

ksi

01 Sep 2005

section of indicated diameter At center

≤13 mm (1=2 in.)

13−25 mm (1=2−1 in.)

At 3=4 radius

At midradius 25−38 mm (1−11=2 in.)

HRC

HB

620−860(a) 90−125

23−30(b )

241−285

1330H 5132H 4130H 8630H

860−1030( 125−150 c)

30−36(d )

285−341

1335H 4135H 50B44H 4042H 8640H 5147H 4047H 94B30H 4137H 5135H 8740H 8645H 3140H 8742H

1030−1170 150−170 (e)

36−41(f)

331−375

1170−1275 170−185 (g)

41−46(h )

46 min(j)

38−50 mm (11=2−2 in.)

50−63 mm (2−21=2 in.)

63−75 mm (21=2−3 in.)

75−89 mm (3−31=2 in.)

89−102 mm (31=2−4 in.)

Oil quenched and tempered

>1275(i)

>185

8737H 50B40H 4140H 8642H 94B40H 94B30H 8740H 3140H

4142H

4142H

4145H

4147H 86B45H 9840H

1340H 5150H 5160H 51B60H 5140H 50B40H 50B50H 8655H 4135H 4137H 4140H 8637H 8642H 94B40H 94B30H 8645H 6150H 3140H 8742H

4145H 4147H 9840H 86B45H 4337H

4150H 4340H

375−429

5145H 5155H 81B45H 86B45H 50B40H 50B44H 4142H 9840H 50B46H 5147H 4145H 4063H 94B40H 8650H 4140H 6150H 8655H 4337H 8640H 8642H 8745H 8740H 8742H

4147H 8660H 4340H

429 min

5150H 5160H 50B60H 5155H 50B50H 51B60H 50B44H 9262H 8660H 5147H 4147H 9260H 8655H 81B45H 8650H 86B45H 6150H

4150H

4150H

4337H 4340H

9850H E4340H

9850H

Water quenched and tempered(k) 620−860(a) 90−125

23−30(b )

241−285

4037H 5130H 5132H 4130H 8630H

5135H 8637H(l)

5140H( m)

4135H 50B40H 8642H 94B30H 3140H

4137H

860−1030( 125−150 c)

30−36(d )

285−341

1330H 5135H

1335H 4135H(l)

1340H( 50B40H 50B44H m) 8640H 5147H 8637H( 8642H 4137H m) 94B30H 8645H 8740H 8742H 3140H

4140H 94B40H

1030−1170 150−170 (e)

36−41(f)

331−375

1330H 1335H 5130H 5132H 5135H 4130H 8620H

4042H 1340H 50B40H (l) 4047H 50B46H 5140H 4137H(l) 4135H 8642H(l) 8637H 94B30H 3140H

8640H( 50B44H 94B40H m) 5147H 8740H( 4140H m) 8645H 8742H

81B45H 4142H 4337H

1170−1275 170−185 (g)

41−46(h )

375−429

5140H 1340H 5145H 50B44H 4037H 50B46H 50B40H (l)


4140H( m)

4142H 81B45H 4337H

4145H 4147H

Page 754



4042H 4137H 8637H

3140H

8640H 5147H(l) 8642H 94B40H (l) 8740H

01 Sep 2005

8645H( m) 8742H( m)

86B45H 9840H 4340H E4340H

81B45H 4147H 5046H 5147H 50B44H (m) 50B46H 4145H 5145H 8645H 4047H 86B45H 4142H 8742H (a) Tensile strength, 790 to 940 MPa (115 to 136 ksi). (b) As-quenched hardness, 42 HRC, or 388 HB. (c) Tensile strength, 940 to 1100 MPa (136 to 160 ksi). (d) As-quenched hardness, 44 HRC, or 415 HB. (e) Tensile strength, 1100 to 1300 MPa (160 to 188 ksi). (f) As-quenched hardness, 48 HRC, or 461 HB. (g) Tensile strength, 1300 to 1530 MPa (188 to 222 ksi). (h) As-quenched hardness, 51 HRC, or 495 HB. (i) Tensile strength, over 1530 MPa (222 ksi). (j) As-quenched hardness, 55 HRC, or 555 HB. (k) Through steels with 0.47% C nominal. (l) May be substituted for steels listed under the 50 to 63 mm (2 to 21=2in.) column at same strength level or less. (m) Not recommended for applications requiring 50% martensite at midradius in sections 38 to 50 mm (11=2to 2 in.) in diameter because of insufficient hardenability. Source: Ref 3 >1275(i)

>185

46 min(j)

429 min

Increasing carbon content consistently increases tensile and yield strength and decreases elongation and reduction in area, regardless of whether the steel is as-rolled or quenched and tempered (provided the ranges of tempering temperatures are the same). However, there is one major disadvantage to increasing the carbon content: Carbon steels show an increasing tendency to crack on quenching as the carbon content increases above about the 0.35% level. Consequently, parts to be made from steel having a carbon content greater than 0.35% should be tested for quench cracking before production is begun. Variations in chemical composition within a specific grade contribute to the scatter of mechanical properties. This is illustrated by the test data in Fig. 10 , where the properties for two heats of quenched and tempered 1050 steel are compared for a tempering range of 315 to 650 °C (600 to 1200 °F). Fig. 10 Effect of composition and tempering temperature on mechanical properties of 1050 steel. Properties are summarized for two heats of 1050 steel that was forged to 38 mm (1.50 in.) in diameter, then water quenched and tempered at various temperatures. Open symbols are for heats containing 0.52 C and 0.93 Mn; closed symbols, for those containing 0.48 C and 0.57 Mn.

Tempering Hardened steels are softened by reheating, although this effect may not be sought in tempering. The real need is to increase the capability of the steel to flow moderately without fracture, and this is inevitably accompanied by a loss of strength. The tensile strength is very closely related to hardness in this class of steels, as heat treated; thus, the effects of tempering can be followed by measuring the Brinell or Rockwell hardness. Figure 11 shows the response to tempering of four carbon and alloy steels containing 0.45% C. All steels were tempered for 1 h at the temperatures indicated. Somewhat shorter or longer intervals at temperature would affect hardness values to various degrees, depending on the tempering temperature.


Page 755



01 Sep 2005

Fig. 11 Tempering characteristics of four 0.45% carbon and alloy steels tempered for 1 h

The general effect of alloying is to retard the tempering rate, and therefore alloy steels require a higher tempering temperature to obtain a given hardness than does carbon steel of the same carbon content. However, the individual elements show significant differences in the magnitude of their retarding effect. Nickel, silicon, aluminum, and, to a large extent, manganese, all of which have little or no tendency to occur in the carbide phase and merely remain dissolved in ferrite, have only a minor effect on the hardness of the tempered steel, as would be expected from the general pattern of solid-solution hardening. However, the carbide-forming elements, chromium, molybdenum, and vanadium, retard softening, particularly at higher tempering temperatures. These elements do not merely raise the tempering temperature; when they are present in higher percentages, the rate of tempering is no longer a continuous function of tempering temperature. That is, the tempering curves for these steels will show a range of tempering temperature in which the tempering is retarded or, with relatively high alloy content, in which the hardness may actually increase with an increase in tempering temperature. This characteristic behavior is known as secondary hardening and results from a delayed precipitation of fine alloy carbides. Secondary hardening is most often encountered in the higher-alloy tool steels. As mentioned previously, the primary purpose of tempering is to impart plasticity or toughness to the steel, and the loss in strength is only incidental to this very important increase in toughness. The increase in toughness after tempering reflects two effects of tempering: • The relief of residual stress induced during quenching • The precipitation, coalescence, and spheroidization of iron and alloy carbides, resulting in a microstructure of greater plasticity In addition to their effects on microstructure, the alloying elements have a secondary function. The higher tempering temperatures for a given hardness, which has been determined to be characteristic of alloy steels (particularly those containing carbide-forming elements), will presumably permit greater relaxation of residual stress and thereby improve properties. Furthermore, as discussed in the section "Alloying Elements in Quenching" in this article, the hardenability of these steels may permit the use of less drastic quenching practices, so that the stress level before tempering will be lower, permitting these steels to be used at a higher level of hardness; this is because higher temperatures are not required for relief of quenching stresses. It should be noted, however, that this latter characteristic is only a secondary function of alloying elements in tempering; the effect primarily reflects the hardenability function of the alloying elements. Another secondary function of alloying elements in tempering is to permit the use of steels with lower carbon content for a given level of hardness, because adequate tempering may be ensured by the retardation of softening caused by alloying. This results in greater freedom from cracking and generally improved plasticity at any given hardness. Here again, the function of alloying elements in tempering is a secondary function; their primary function is to increase hardenability sufficiently to offset the effect of a decreased carbon content. The increase in plasticity upon tempering is discontinuous in those alloy steels that contain the carbide-forming elements; the behavior of notched specimens shows a characteristic irregularity at approximately 260 to 315 °C (500 to 600 °F). The quenched martensitic steel gains toughness, as reflected in a notched-bar impact test, by tempering at temperatures as high as 205 °C (400 °F). However, after tempering at higher temperatures, in the temper-brittle range, these types of steel lose toughness until they may be less tough than the same steels not tempered. Still higher tempering temperatures restore greater toughness (see Fig. 12 ). Fig. 12 Hardness and notch toughness of 4140 steel tempered for 1 h at various temperatures


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The mechanism of this behavior is not fully understood, but it seems to be associated with the first precipitation of carbide particles and is presumably a grain boundary phenomenon; fractures of steels tempered in this region tend to follow intergranular paths. Thus, there is a range of tempering temperatures at about 205 to 370 °C (400 to 700 °F) never used for these steels; the tempering temperature is either below 205 °C (400 °F) or above 370 °C (700 °F). Although this phenomenon is common to all of these alloy steels, the alloying elements have a secondary function in this connection; a combination of carbon and alloy contents of suitable hardenability may be chosen that would permit tempering to the desired strength at temperatures outside this undesirable range. Temper brittleness is another example of a discontinuous increase in plasticity subsequent to the tempering of steels containing the carbide-forming elements. This phenomenonis manifested as a loss of toughness, observed after slow cooling from tempering temperatures of 575 °C (1070 °F) or higher or after tempering in the temperature range between approximately 375 and 575 °C (700 and 1070 °F). Thus, a steel that is susceptible to temper embrittlement may lose much of its plasticity, as indicated by a notched-bar impact test, during slow cooling from a high tempering temperature, although the same steel will be very tough if it is quenched from the same tempering temperature. This expedient of quenching from the tempering temperature is often overlooked as a practical means for avoiding sever temper embrittlement in susceptible steels tempered at 575 °C (1070 °F) or higher. In steels susceptible to temper brittleness, embrittlement will also be observed after tempering at 375 to 575 °C (700 to 1070 °F), particularly if the tempering times are protracted. Under such circumstances, quenching from the tempering temperature will never restore the toughness. High manganese, phosphorus, and chromium concentrations appear to accentuate the embrittling reaction; molybdenum has a definite retarding effect. Here again, the carbon and alloying elements may be chosen so that the susceptibility to temper embrittlement is minimized or the desired strength level is obtained by tempering either below 375 °C (700 °F) or above 575 °C (1070 °F) and then quenching. Temper brittleness is discussed in greater detail in the article"Embrittlement of Steels" in this Volume.

Distortion in Heat Treatment Distortion during heat treatment may occur with almost any hardenable carbon or alloy steel, although distortion is usually more severe for carbon grades than for alloy grades of equivalent carbon content. Carbon steels distort more than alloy steels mainly because carbon steels require a water or brine quench to develop full hardness (at least in sections thicker than about 9.5 mm, or 3=8in.). This often eliminates carbon steels from consideration for critical parts. This distortion may be observed as a change in dimensions (size distortion) or a change in configuration or contour (shape distortion or warpage), or both. A more complete discussion of these types of distortion and the factors that influence them may be found in Ref 4, 5, and 6. Several factors contribute to the total distortion that occurs during heat treatment. These include residual stresses that may be present as a result of machining or other cold-working operations, the method of placing in the furnace, the rate of heating, nonuniform heating, and the normal volumetric changes that occur with phase transformations. However, the most important, single factor is uneven cooling during quenching, caused mainly by the configuration and by changes in cross-sectional area. Symmetrical parts with little or no variation in section may have almost no distortion, whereas complex parts with wide variations in section may distort so much that they cannot be used (or at least so much that they require excessive finishing operations to make them suitable for use). Other factors being equal, the distortion in carbon steels will increase as the carbon content increases because of the gradual lowering of the martensite start (Ms) temperature with increasing carbon. There is also a significant variation in the magnitude of distortion and direction of dimensional change among different heats of the same grade of steel, even though other variables are minimal. This happens because of several factors, including minor variations in composition and grain size, but mainly because of the history of the steel with regard to hot working, cold working, and heat treatment. Because of the different variables that contribute to total distortion, the prediction of distortion in actual parts is seldom reliable if it is based on the behavior of small test pieces. The most practical approach is to make studies on pilot lots of actual pieces that have been heat treated under production conditions. This procedure eliminates the shape variable so that the direction and magnitude of distortion can be plotted as ranges that incorporate most of the other variables. After a quantity of such data has


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been secured, a series of guideposts is established, and it becomes possible to predict distortion for similar parts made from the same steel grade with reasonable accuracy. However, it must be emphasized that any such study is accurate only when many parts made from different heats supplied by several mills are included.

Induction and Flame Hardening The relatively low hardenability of carbon steels is often a primary reason for choosing them in preference to alloy steels for parts that are to be locally heat treated by flame or induction hardening. One of the oldest rules for selecting steels for heat treating is to choose grades that are no higher in carbon or alloy content than is essential to develop required properties. This rule remains valid in the selection of steels to be heat treated by induction or flame processes. When the peripheries of steel parts are heated rapidly and quenched, the tendency to crack depends mainly on a combination of four factors: • • • •

Final surface hardness Temperature to which the surface has been heated Uniformity of heating Depth of hardened zone

The optimum heat pattern for either induction or flame heating depends on the type of steel and on the mass and shape of the part. The ideal heat pattern for any specific part will provide a hardened shell to a depth that will strengthen the part by establishing a favorable stress pattern. However, if the hardened zone is too deep for the specific section thickness, high tensile stresses are established in the surface layers, and these may either cause cracking or adversely affect service life. Excessive depth of the hardened zone can be caused by improper processing (overheating, for instance) or by the choice of a steel with excessive hardenability. However, excessive carbon can aggravate other contributing factors and become the basic cause for cracking. The Ms temperature decreases as the carbon content increases. It is lowered further by higher austenitizing temperatures. In general, as the Ms temperature is lowered, the probability of surface cracking increases.

Fabrication of Parts and Assemblies Fabrication processes are usually performed on hardenable carbon and alloy steels in the unhardened condition, that is, prior to heat treating. This is done primarily to avoid the high cost and difficulty of fabrication that are characteristic of high-strength materials. However, even in the unhardened condition, there are differences among the various grades in respect to formability, weldability, machinability, and forgeability properties. In many instances, difficulties arising during the fabrication of a given hardenable steel are directly related to the maximum hardness that can be developed and to hardenability. REFERENCES 1. 2. 3. 4. 5. 6.

1989 SAE Handbook, Vol 1, Materials, Society of Automotive Engineers, 1989 S.L. Semiatin and D.E. Stutz, Induction Heat Treatment of Steel, American Society for Metals, 1986, p 24 Republic Alloy Steels, Republic Steel Corporation, 1961 B.S. Lement, Distortion in Tool Steels, American Society for Metals, 1959 Properties and Selection of Tool Materials, American Society for Metals, 1975 J.A. Ferrante, Controlling Part Dimensions During Fabrication and Heat Treatment, Met. Prog., Vol 87 (No. 1), Jan 1965, p 87−90; reprinted in Source Book on Heat Treating, Vol I, American Society for Metals, 1975


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Hardenability of Carbon and Low-Alloy St...

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ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Hardenability of Carbon and Low-Alloy Steels Hardenability of Carbon and Low-Alloy Steels Revised by Harold Burrier, Jr., The Timken Company HARDENABILITY OF STEEL is the property that determines the depth and distribution of hardness induced by quenching. Steels that exhibit deep hardness penetration are considered to have high hardenability, while those that exhibit shallow hardness penetration are of low hardenability. Because the primary objective in quenching is to obtain satisfactory hardening to some desired depth, it follows that hardenability is usually the single most important factor in the selection of steel for heat-treated parts. Hardenability should not be confused with hardness as such or with maximum hardness. The maximum attainable hardness of any steel depends solely on carbon content. Also, the maximum hardness values that can be obtained with small test specimens under the fastest cooling rates of water quenching are nearly always higher than those developed under production heat-treating conditions, because hardenability limitations in quenching larger sizes may result in less than 100% martensite formation. The effects of carbon and martensite content on hardness are shown in Fig. 1 . Basically, the units of hardenability are those of cooling rate, for example, degrees per second. These cooling rates, as related to the continuous-cooling-transformation behavior of the steel, determine the hardness and microstructural outcome of a quench. In practice, these cooling rates are often expressed as a distance, with other factors such as the thermal conductivity of steel and the rate of surface heat removal being held constant. Therefore, the terms Jominy distance and ideal critical diameter can be used. Fig. 1 Effect of carbon on the hardness of martensite structures


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The hardenability of steel is governed almost entirely by the chemical composition (carbon and alloy content) at the austenitizing temperature and the austenite grain size at the moment of quenching. In some cases, the chemical composition of the austenite may not be the same as that determined by chemical analysis, because some carbide may be undissolved at the austenitizing temperature. Such carbides would be reflected in the chemical analysis, but because the carbides are undissolved in the austenite, neither their carbon nor alloy content can contribute to hardenability. In addition, by nucleating transformation products, undissolved carbides can actively decrease hardenability. This is especially important in high-carbon (0.50 to 1.10%) and alloy carburizing steels, which may contain excess carbides at the austenitizing temperature. Consequently, such factors as austenitizing temperature, time at temperature, and prior microstructure are sometimes very important variables when determining the basic hardenability of a specific steel composition. Certain ingot casting and hot reduction practices may also develop localized or periodic inhomogeneities within a given heat, further complicating hardenability measurements. The effects of all these variables are discussed in this article.

Hardenability Testing The hardenability of a steel is best assessed by studying the hardening response of the steel to cooling in a standardized configuration in which a variety of cooling rates can be easily and consistently reproduced from one test to another. The Jominy end-quench test fulfills the cooling rate requirements of hardenability testing of a broad range of alloy steels. The test specimen, a 25.4 mm (1.000 in.) diam bar 102 mm (4 in.) in length, is water quenched on one end face. The bar from which the specimen is made must be normalized before the test specimen in machined. The test involves heating the test specimen to the proper austenitizing temperature and then transferring it to a quenching fixture so designed that the specimen is held vertically 12.7 mm (0.5 in.) above an opening through which a column of water can be directed against the bottom face of the specimen (Fig. 2 a). While the bottom end is being quenched by the column of water, the opposite end is cooling slowly in air, and intermediate positions along the specimen are cooling at intermediate rates. After the specimen has been quenched, parallel flats 180° apart are ground 0.38 mm (0.015 in.) deep on the cylindrical surface. Rockwell C hardness is measured at


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intervals of 1=16in. (1.6 mm) for alloy steels and 1=32in. (0.8 mm) for carbon steels, starting from the water-quenched end. A typical plot of these hardness values and their positions on the test bar, as shown in Fig. 2 (b), indicates the relation between hardness and cooling rate, which in effect is the hardenability of the steel. Figure 2 (b) also shows the cooling rate for the designated test positions. Details of the standard test method are available in ASTM A 255 and SAE J406. Fig. 2 Jominy end-quench apparatus (a) and method for presenting end-quench hardenability data (b)

The Carburized Hardenability Test. It is often necessary to determine the hardenability of the high-carbon case regions of carburized steels. Such information is important in controlling carburizing and quenching practice and in determining the ability of a specific steel to meet the microstructural and case depth requirements of the carburized component manufactured from the steel. As a general rule, adequate core hardenability does not ensure adequate case hardenability, especially when it is required to reheat for hardening after carburizing rather than to quench directly from the carburizing furnace. Two factors are responsible for this fact. The first is that equal alloying additions do not have the same effect on the hardenability of all carbon levels of alloyed steels. The second factor (as noted earlier) is that the high-carbon case regions do not always achieve full solution of alloy and carbides, as is normally achieved in the austenite of the low-carbon core region, prior to quenching. Accordingly, direct measurements of case hardenability are very important whenever a carburizing steel must be selected for a specific application. Measurements of case hardenability are performed as follows. A standard end-quench bar is pack carburized for 9 h at 925 °C (1700 °F) and end quenched in the usual manner. A comparison bar is simultaneously carburized in the same pack to determine carbon penetration. Successive layers are removed from it and analyzed chemically to determine the carbon content at various depths. When a carbon-penetration curve is established, depths to various carbon levels can be determined in the Jominy bar, assuming that the distribution of carbon in the end-quench specimen is the same as in the carbon gradient bar. Longitudinal flats are then carefully ground to various depths on the end-quench bar (usually to carbon concentrations of 1.1, 1.0, 0.9, or 0.8%, and in some cases to as low as 0.6%), and hardenability is determined at these carbon levels by hardness traverses. In grinding, care must be exercised to avoid overheating and tempering, and in conducting hardness surveys, similar concern must be shown to ensure that the hardness level corresponds to a single carbon level by remaining in the exact center of the flat. Rockwell A hardness readings are preferable to Rockwell C readings because they minimize the depth of indentor penetration into softer subsurface layers. Rockwell A values are converted into Rockwell C values for plotting, as illustrated in Fig. 3 , which shows the curves of carburized hardenability of an EX19 steel. In the higher-carbon layers of carburized specimens, the hardness will be influenced by the presence of retained austenite. Therefore, it is often useful to evaluate the microstructure/depth relationship by metallographically polishing and etching the ground flats. The Jominy distance to some chosen level of nonmartensitic transformation product can then be used as a measure of hardenability. Fig. 3 Carburized hardenability, EX19 steel. Composition: 0.18 to 0.23% C, 0.90 to 1.20% Mn, 0.40 to 0.60% Cr, 0.08 to 0.15% Mo, 0.0005% B (min)


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The case hardenability of steels that are carburized and then reheated for hardening at temperatures below 925 °C (1700 °F), such as 8620, 4817, and 9310, can also be determined by using a modification of this technique. The carburized end-quench specimens and companion gradient bars are oil quenched together from carburizing, but are then reheated in an atmosphere furnace to the desired austenitizing temperature for a total of 55 to 60 min, which should ensure at least 30 to 35 min at temperature. The hardenability specimen is then end quenched, and the carbon gradient bar is oil quenched and tempered to facilitate machining for carbon gradient determination, as described above. It is recommended that case hardenability tests be performed on no fewer than two test specimens. A more detailed description of the case hardenability measurement technique appears in SAE J406. Air Hardenability Test. Occasionally, the hardening performance either of a steel cooled at a rate slower than that applied to the end-quench bar or of steels of very high hardenability must be determined. An air hardenability test method described in Ref 1 can be employed for this purpose. In this test, a machined and partially threaded round test specimen, 25.4 mm (1.000 in.) in diameter and 254 mm (10 in.) long, is inserted to a depth of 152 mm (6 in.) in a hole drilled in a bar 152 mm (6 in.) in diameter and 381 mm (15 in.) long, thus leaving 102 mm (4 in.) of the test bar length exposed (Fig. 4 ). A second test specimen can be inserted at the opposite end of the bar holder to serve as a duplicate. With both test bars securely in place, the assembly is heated to the proper austenitizing temperature, after which it is transferred to a convenient location for cooling in still air. This cooling procedure results in very slow and ever decreasing cooling rates along the length of the test bars. Hardness is then measured at discrete intervals along each test bar and plotted against distance from the exposed end on charts specifically designed for this purpose. Fig. 4 Dimensions (given in inches) of components in air hardenability test setup

Continuous-Cooling-Transformation Diagrams. The use of continuous-cooling-transformation diagrams determined


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dilatometrically, for example, can also be helpful in evaluating the cooling behavior of high-hardenability steels.

Low-Hardenability Steels In plain carbon and very low-alloy steels, the cooling rate at even the 1.6 mm (1=16in.) position on a standard Jominy bar may not be fast enough to produce full hardening. Therefore, this test lacks discrimination between these steels. Tests that are more suited to very low hardenability steels include the hot-brine test and the surface-area-center (SAC) test. In the hot-brine test proposed by Grange, coupons (Fig. 5 ) are quenched in brine maintained at a series of different temperatures. As shown in Fig. 6 , the resulting hardnesses provide a very sensitive test of hardenability. Fig. 5 Hot-brine hardenability test specimen. (a) Specimen dimensions. (b) Method of locating hardness impressions after heat treatment. Dimensions given in millimeters. Source: Ref 2

Fig. 6 Typical results of the hot-brine hardenability test. Steel composition: 0.18% C, 0.81% Mn, 0.17% Si, and 1.08% Ni. Austenitized at 845 °C (1550 °F). Grain size: 5 to 7. RT, room temperature. Source: Ref 2

In the SAC test, a 25.4 mm (1.000 in.) round bar is normalized by cooling in air and then reaustenitized for water quenching. Hardnesses are measured on a specimen cut from the center of the 100 mm (4 in.) length. Hardness is determined on the surface, the center, and at 1.6 mm (1=16in.) intervals from surface to center. An area hardness is then computed as the sum of


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the average hardness in each interval × 1=16(Fig. 7 ). The resulting set of three-digit numbers, for example, SAC No. 63-52-42, indicates a surface hardness of 63 HRC, a Rockwell-inch area of 52, and a center hardness of 42 HRC. Testing details are given in SAE J406. Fig. 7 Surface-area-center estimation of area

Calculation of Hardenability The hardenability of a steel is primarily a function of the composition (carbon, alloying elements, and residuals) and the grain size of the austenite at the instant of quenching. If this relationship can be determined quantitatively, it should be possible to predict the hardenability of a steel through a relatively simple calculation. Such a technique was published by Grossmann in 1942, based on his observation that hardenability could be expressed as the product of a series of composition-related multiplying factors (Ref 3). The result of the calculation is an estimate of DI, the ideal critical diameter of the steel. The multiplying-factor principle is still used today in several hardenability calculation techniques (see the article "High-Strength Structural and High-Strength Low-Alloy Steels" in this Volume for examples of multiplying factors for quench and tempered low-alloy steels). Other researchers have developed methods based on regression equations and on calculation from thermodynamic and kinetic first principles. To date, none of the hardenability prediction methods has proved to be universally applicable to all steel types; that is, different predictors are more suited to steels of given alloying systems, carbon contents, and hardenability levels. In addition, it is often necessary to fine-tune the predictions based on the characteristics (residuals, melt practice, and so on) of a particular steel producer. Some excellent discussions of current thinking on this subject are available in Ref 4 and 5. Properly used, hardenability calculations can provide a valuable tool for designing cost-effective alternative steels, for deciding the disposition of heats in the mill prior to rolling, and possibly for replacing the costly and time-consuming measurement of hardenability.

Effect of Carbon Content Carbon has a dual effect in hardenable alloy steels: It controls maximum attainable hardness and contributes substantially to hardenability. The latter effect is enhanced by the quality and type of alloying elements present. It might be concluded, therefore, that increasing the carbon content is the least expensive approach to improving hardenability. This is true to a degree, but several factors weigh against the use of large amounts of carbon: • High carbon content generally decreases toughness at room and subzero temperatures • It produces harder and more abrasive microstructures in the annealed conditions, which makes cold shearing, sawing,


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machining, and other forms of cold processing more difficult • It makes the steel more susceptible to hot shortness in hot working • It makes the steel more prone to cracking and distortion in heat treatment. Because of these disadvantages, more than 0.60% C is seldom used in steels for machine parts, except for springs and bearings, and steels with 0.50 to 0.60% C are used less frequently than those containing less than 0.50% C Figure 8 shows the differences between minimum hardenability curves for six series of steels. In each series, alloy content is essentially constant, and the effect of carbon content on hardenability can be observed over a range from 0.15 to 0.60%. The hardness effect is shown by the vertical distance between the curves at any position on the end-quench specimen, that is, for any cooling rate. This effect varies significantly, depending on the type and amounts of alloying elements. For example, referring to Fig. 8 (d) to (f), an increase in carbon content from 0.35 to 0.50% in each of the three series of steels causes hardness increases (in Rockwell C points) at four different end-quench positions, as shown below:

Series

1

=16

Distance from quenched surface, in. 4 8 =16 =16

12

=16

41xxH

8

10

17

20

51xxH

8

13

9

8

86xxH

8

12

18

12

Fig. 8 Effect of carbon content on the minimum end-quench hardenability of six series of alloy H-steels. The number adjacent to each curve indicates the carbon content of the steel, to be inserted in place of xx in alloy designation.

The hardenability effect of carbon content is read on the horizontal axis in Fig. 8 . If the inflection points of the curves are used to approximate the position of 50% martensite transformation, the effect of carbon content on hardenability in 8650 versus 8630 steel can be expressed as +4=16; that is, the inflection point is moved from the 5=16position to the 9=16position. Similarly, with nominal carbon contents of 0.35 and 0.50%, the hardenability effect of carbon is seen to be less (2=16) in 51xx series steels and more (6=16) in 41xx steels. Considering the combined hardening and hardenability effects in terms of quenching speed, the cooling rate (or quenching speed) required to produce 45 HRC is affected more by 0.15% C with certain combinations of alloying elements than it is by


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other combinations. For example, in a steel containing 0.75 Cr and 0.15 Mo (a 41xxH series steel, for example), increasing the carbon content by 0.15% lowers the required or critical cooling rate to obtain 45 HRC from 25 to 4.6 °C (45 to 8.3 °F) per second, while in a steel containing 0.75% Cr and no molybdenum (51xxH series), the same increase in carbon content lowers the cooling rate from 47 to 21 °C (85 to 37 °F) per second. The practical significance of the effect of carbon and alloy contents on cooling rate is considerable. In a 51 mm (2 in.) diam bar of 4150 steel, a hardness of 45 HRC can be obtained at half-radius using an oil quench without agitation. In a 4135 steel bar of the same diameter, to obtain the same hardness at half-radius would require a strongly agitated water quench. Comparing 32 mm (11=4in.) diam bars of 5135 and 5150 steel, an agitated water quench will produce a hardness of 45 HRC at half-radius in the 5135 bar; the identical condition can be obtained in the 5150 bar using an oil quench with moderate agitation. Thus, an increase or decrease in carbon content or an alloying addition, such as 0.15% Mo, affects the results obtained both in terms of the quenching severity required and the section size in which the desired results can be obtained. Figure 9 shows how steels are rated on the basis of ideal critical diameter by expressing the effect of carbon and alloy content on the section size that will harden to 50% martensite at the center, assuming an ideal quench. An ideal quench is defined as one that removes heat from the surface of the steel as fast as it is delivered to the surface. In general, the relation between hardness and carbon content that is important in practice is obscured in this rating method because the steel is rated to a constant microstructure. Hardness decreases continuously with lower carbon contents. Fig. 9 Effect of carbon content on ideal critical diameter, calculated for the minimum chemical composition of each grade

Alloying Elements The most important function of the alloying elements in heat-treatable steel is to increase hardenability. Increased hardenability makes possible the hardening of larger sections and the use of an oil rather than a water quench to minimize distortion and to avoid quench cracking. When the standard alloy steels are considered, it is found that, for practical purposes, all compositions develop the same tensile properties when quenched to martensite and tempered to the same hardness below 50 HRC. However, it should not be inferred that all tempered martensites of the same hardness are alike in all respects. For example, plain carbon martensites have lower reduction-in-area values than alloy martensites. A further difference, sometimes important, is that fully quenched alloy steels require, for the same hardness levels, higher tempering temperatures than carbon steels. This difference in tempering temperature may serve to reduce the residual stress level in finished parts. The stress reduction could be an advantage or a disadvantage, depending on whether a controlled compressive stress is desired in the part. Although tensile properties may not differ significantly from one alloy steel to another, considerable differences may exist in fracture toughness and low-temperature impact properties. In general, steels with a higher nickel content, such as 4320, 3310, and 4340, offer much greater toughness at a given hardness level. In some applications, the toughness factor rather than hardenability may dictate steel selection, but hardenability is still important, because steels that can be fully quenched to 100% martensite are much tougher than those that cannot. Usually, the least expensive means of increasing hardenability at a given carbon content is by increasing the manganese content. Chromium and molybdenum, already referred to as increasing hardenability, are also among the most economical elements per unit of increased hardenability. Nickel is the most expensive per unit, but is warranted when toughness is a primary consideration.


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Important synergistic effects, not yet fully defined, can also occur when combinations of alloying elements are used in place of single elements. Some examples of known synergistic combinations are nickel plus manganese, molybdenum plus nickel, and silicon plus manganese. Boron. Another potent and economical alloying element is boron, which markedly increases hardenability when added to a fully deoxidized steel. The effects of boron on hardenability are unique in several respects: • • • •

A very small amount of boron (about 0.001%) has a powerful effect on hardenability The effect of boron on hardenability is much less in high-carbon than in low-carbon steels Nitrogen and deoxidizers influence the effectiveness of boron High-temperature treatment reduces the hardenability effect of boron

Recommended austenitizing temperatures for boron H-steels are given with the H-bands. Figure 10 illustrates the very small amount of boron required for an optimum hardenability effect when appropriate protection of the boron is afforded by additions of titanium or zirconium. In carburizing steels, the effect of boron on case hardenability may be completely lost if nitrogen is abundant in the carburizing atmosphere. The cost of boron is usually much less than that of other alloying elements having approximately the same hardenability effect. Fig. 10 Influence of effective boron content (βeff) on the hardenability of an 8620 type steel. βeff = B-[(N-0.002)-Ti/5-Zr/15] » 0. Source: Ref 5

Effect of Grain Size The hardenability of a carbon steel may increase as much as 50% with an increase in austenite grain size from ASTM 8 (6 to 10) to ASTM 3 (1 to 4). The effect becomes more pronounced if the carbon content is increased at the same time. When the danger of quench cracking is remote (no abrupt changes in section thickness) and engineering considerations permit, it may sometimes appear to be more practical to use a coarser-grain steel rather than a fine-grain or more expensive alloy steel to obtain hardenability. However, this is not recommended, because the use of coarser-grain steels usually involves a serious sacrifice in notch toughness and may lead to other difficulties.

Variations Within Heats Segregation of carbon, manganese, and other elements always occurs during ingot pouring and solidification. As a result, the hardenability of the steel in these segregated portions will differ from that in the remainder of the ingot. In general, specimens taken from the top of the ingot have higher hardenability than steel from the middle, and specimens from the bottom of the ingot will have lower hardenability than steel from the middle. This gradual increase in hardenability from the bottom of the first ingot to the top of the last ingot is illustrated in Fig. 11 (a) for 16 heats of 1035 carbon steel. The hardenability spread for 8 heats of 1035 steel containing 0.05 to 0.12% Mo, plotted in Fig. 11 (b), shows a similar trend. Comparison between Fig. 11 (a) and 11 (b) shows the effect of molybdenum on hardenability. Fig. 11 Effect of test location of (a) 1035 steel and (b) 1035 steel with 0.05 to 0.12% Mo on SAC (Rockwell-in.) hardenability


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The same effect is observed in alloy steels. End-quench hardenability test results for one heat of 4028 steel (Fig. 12 ) show higher hardenability for a cast bar taken from the top of the last ingot of the heat than for a specimen from the melting floor and labeled cast end-quench specimen. The latter was taken from about the middle of the heat. After the heat of steel was rolled, the hardenability was slightly lower, as shown by the curve representing results on eight end-quench specimens. Data for 465 heats of ten other steels are summarized in Fig. 13 . Fig. 12 Variation of hardenability within a heat of 4028 steel

Fig. 13 Variation in hardenability from first to last ingot in heat for several carbon and alloy steels


Page 768



01 Sep 2005

Effect of Hot Working. Processing variables, such as the amount of hot working and the location of the test specimen in the semifinished section, have an effect on hardenability. A 330 mm (13 in.) square bloom of 1330 steel was forged progressively to bar sizes of 305, 255, 205, and 150 mm (12, 10, 8, and 6 in.) in diameter. Each bar size was evaluated by tests on end-quench specimens cut from five locations (center, quarter-radius, half-radius, three-quarter-radius, and just below the surface). Data in Fig. 14 show that the variation in hardenability narrows as the bar size is decreased by hot work. Fig. 14 Effect of hot working and location of test bars on end-quench hardenability of 1330 steel. A 330 mm (13 in.) bloom was progressively forged to bars of the diameters shown.


Page 769



01 Sep 2005

Determining Hardenability Requirements The basic information needed to specify a steel with adequate hardenability includes: • The as-quenched hardness required prior to tempering to final hardness that will produce the best stress-resisting microstructure • The depth below the surface to which this hardness must extend • The quenching medium that should be used in hardening As-Quenched Hardness. The Iron and Steel Technical Committee of the Society of Automotive Engineers (SAE) War Engineering Board approved and issued the relation shown in Fig. 15 (a) as a recommendation for as-quenched hardness as a function of the hardness desired after tempering. Figure 15 (a) does not specify the degree of hardening (percentage martensite) preferred in obtaining the as-quenched hardnesses indicated. It is possible, as shown in Fig. 15 (b), to select steels that will produce these hardnesses with less than 90% martensite.


Page 770



01 Sep 2005

Fig. 15 Curves for steel selection based on hardness. (a) Minimum as-quenched hardness to produce various final hardnesses after tempering. (b) Dependence of as-quenched hardness on percentages of martensite and carbon

To ensure optimum properties, common practice is to select the steel with the lowest carbon content that will produce the indicated as-quenched hardness using the quenching medium available (or one that can be made available). Following this procedure, the structures possessing the indicated hardnesses would be fully hardened; that is, they would contain more than 90% martensite, which is a common and practical definition of full hardening and the one employed by the SAE committee. For components subjected to bending in service, it is considered adequate to have 90% martensite at the three-quarter-radius location. To ensure this, hardness levels are specified at half-radius. Depth of Hardening. The depth and percentage of martensite to which parts are hardened may affect their serviceability, but it always affects the hardenability required and therefore the cost. In parts less highly stressed in bending, hardening to 80% martensite at three-quarter-radius of the part as finished may be sufficient; in other parts, even less depth may be required. the latter include principally those parts designed for low deflection under load, in which even the exterior regions are only moderately stressed. In contrast, some parts loaded principally in tension and others operating at high hardness levels, such as springs of all types, are usually hardened more nearly through the section. In automobile leaf springs, the leaves are designed with a low section modulus in the direction of loading. The allowable deflection is large, and most of the cross section is highly stressed. In general, hardening need be no deeper than is required to provide the strength to sustain the load at a given depth below the surface. Therefore, parts designed to resist only surface wear, pure bending, or rolling contact often do not justify the cost of providing the hardenability required for hardening through the entire cross section. When service requirements mandate that hardening must produce more than 80% martensite, the section size that can be hardened to a prescribed depth decreases rapidly as the percentage of martensite required increases. For example, let us assume that 95% martensite (51 HRC minimum hardness) is required in 8640H steel. Then the largest section size that can be hardened to the center in oil would be 16 mm (5=8in.); a 25 mm (1 in.) section could be hardened to only three-quarter-radius. Again, on the basis of 95% martensite, the deepest hardening of standard steels, 4340H, will harden to the center of a 51 mm (2 in.) section; on the basis of 80% martensite (45 HRC), a 92 mm (35=8in.) round will harden to the center in oil. The above examples emphasize the need for engineering judgment in requiring very deep hardening or unusually high percentages of martensite. When these requirements are not wholly justified, the results are overspecification of steel at higher cost and greater likelihood of distortion and quench cracking. Quenching Media. The cooling potential of quenching media is a critical factor in heat-treating processes because of its contribution to attaining the minimum hardenability requirement of the part or section being heat treated. The cooling potential, a measure of quenching severity, can be varied over a rather wide range by: • Selection of a particular quenching medium • Control of agitation • Additives that improve the cooling capability of the quenchant Any or all of these variables can be employed to increase quenching severity and provide the following advantages: • • • •

Permit the use of less expensive (lower-alloy) steels of lower hardenability Optimize the properties of the steel selected Permit the use of less expensive quenching media Improve productivity and achieve cost reductions as a result of shorter cycle times and higher production rates

In practice, however, two other considerations modify the selection of quenching medium and quenching severity: the amount of distortion that can be tolerated and the susceptibility to quench cracking.


Page 771



01 Sep 2005

In general, the more severe the quenchant and the less symmetrical the part being quenched, the greater the size and shape changes that result from quenching and the greater the risk of quench cracking. Consequently, although water quenching is less costly than oil quenching and water-quenched steels are less expensive than those requiring oil quenching, it is important that the parts to be hardened be carefully reviewed to determine whether the amount of distortion and the possibility of cracking as a result of water quenching will permit taking advantage of the lower cost of water quenching. Oil, salt, and synthetic water-polymer quenchants are alternatives, but their use often requires steels of higher alloy content to satisfy hardenability requirements. A rule regarding selection of a steel and quenching medium for a given part is that the steel should have a minimum hardenability not exceeding that required by the quenching severity of the medium selected. The steel should also contain the lowest carbon content compatible with the required hardness and strength properties. This rule is based on the fact that the quench cracking susceptibility of steels increases with a decrease in Ms temperature and/or an increase in carbon content. Table 1 lists typical quenching severity, or H, values for the common quenching media and conditions. These data are for media containing no additives. Figure 16 shows the effects of additives and of other quenching media. According to these data, considerable improvement in the cooling capability of quenchants can be obtained by such additions as water to hot salt, proprietary additives to oil, and polyalkylene glycol (polymer) to water. The polymer-water mixtures polyacrylamide gel (PAG), polyvinyl pyrrolidone (PVP), and polyvinyl alcohol (PVA) are gaining favor because they can be made to span the quenching severity range from oil to water by simple variation of the glycol (polymer) concentration in water. Also, because they are free of fire hazards and obnoxious environmental pollution agents, they have no adverse effect on working conditions. The quenching severity of these media should be tested at frequent intervals because dragout and thermal breakdown may affect their quenching efficiency. Table 1 Quenching severities, H, for various media and quenching conditions Typical flow rates Quenchant agitation

Typical H values

m/min

sfm

Air

Mineral oil

Water

Brine

None

0

0

0.02

0.20−0.30

0.9−1.0

2.0

Mild

15

50

...

0.20−0.35

1.0−1.1

2.1

Moderate

30

100

...

0.35−0.40

1.2−1.3

...

Good

61

200

0.05

0.40−0.60

1.4−2.0

...

Strong

230

750

...

0.60−0.80

1.6−2.0

4.0

Fig. 16 Approximate quench severities for quenching media containing additives to improve cooling capacity

Hardenability Versus Size and Shape. When end-quench data such as those shown in Fig. 2 are available, either of two methods can be used to estimate the hardenability a steel part of given size and configuration must have to achieve the desired hardness, strength, and microstructure at critical locations when quenched in various production media. These methods are: • Method 1: The correlation of end-quench hardness data (Jeh) with equivalent hardness locations in variously quenched shapes


Page 772



01 Sep 2005

• Method 2: The correlation of end-quench cooling rate data (Jec) with equivalent cooling rate locations in variously quenched production shapes. Method 1 (Fig. 17 ) is the more accurate and preferred method, because in practice it has been found that, when cooling at the same rates, large sections produce somewhat lower hardnesses than smaller sections, including end-quench and air hardenability bars. This difference has been attributed to two factors (Ref 6): • Higher contraction stresses in large parts accentuate the transformation of austenite • Quenching severity, H, decreases with an increase in section size Also, in using the cooling rate method (method 2), it is difficult to determine cooling rates with a high degree of accuracy. Nevertheless, correlations that equate cooling conditions along the end-quench bar (J ec) with those in production shapes quenched in various liquid media are also extremely useful when attempting to establish the required hardenability and/or quenching conditions for a production part. Fig. 17 Determination of Jominy equivalent hardness (Jeh) rates

Jominy equivalent hardness (Jeh) rates are determined by comparing the hardnesses of cross sections of parts receiving the established production heat treatment to hardnesses obtained on end-quenched bars of the same steel. A typical procedure is as follows: 1. Select hardening and quenching conditions that the production hardening equipment can easily fulfill. 2. Select a low-hardenability steel, such as 8620, 4023, or 1040, and manufacture a quantity of finished components: gears, bearings, shafts. 3. Quench a number of these components (in the uncarburized condition) in the production facility. 4. Measure the hardnesses obtained at all critical locations from the surface to the core. 5. Compare the measured hardness values at these locations with equivalent hardness values produced at some end-quench (Jeh) location on a Jominy bar made from the same heat and end quenched from the same thermal conditions. 6. The Jeh values obtained in this fashion define the equal hardness cooling conditions for each location in the production-quenched component. 7. Finally, select from available end-quench data a steel that will produce the hardnesses required at each critical Jeh location in the finished production part. If end-quench data are not available, calculate a suitable composition by one of the standard methods.

General Hardenability Selection Charts


Page 773



01 Sep 2005

Figures 18 and 19 show correlations of Jec equivalent cooling rates in end-quench hardenability specimens and round bars of up to 102 mm (4 in.) in diameter when quenched in oil, water, brine, and hot salt at various controlled agitation rates. They correlate bar diameter with equivalent positions on the end-quench hardenability specimen for ten modes of quenching, for both scaled and scale-free bars, and with data grouped according to bar location instead of by quenching mode. Fig. 18 Correlation of Jec equivalent cooling rates in the end-quench hardenability specimen and round bars quenched in oil, water, and brine, (a), (c), and (e) Nonscaling austenitizing atmosphere. (b), (d), and (f) Austenitized in air

Brine (1), violent agitation 4.0

Water (2), 60 m/min (200 sfm)

Still water (3)

Oil (4), 230 m/min (750 sfm)

Oil (5), 60 m/min (200 sfm)

Oil (6), 15 m/min (50 sfm)

Still oil (7)

1.5

1.0

0.8

0.5

0.35

0.20

Fig. 19 Correlation of Jecequivalent cooling rates in the end-quenched hardenability specimen and round bars quenched in salt at 205 °C (400 °F)


Page 774



01 Sep 2005

Table 2 has been devised to work with the charts in Fig. 18 and 19 and includes most of the steels for which H-bands have been established, showing the location on the end-quenched specimen of the low limit of the H-band for six different hardness levels that might be specified for as-quenched hardness; 55, 50, 45, 40, 35, and 30 HRC. The last two levels apply primarily to the core hardness of carburized parts. Table 2 Classification of H steels according to minimum hardness at various distances from quenched end

Distance from quenched end, 1 =16 th in.

H steels with a minimum Typical values(a) obtained by the use of Fig. 18 bar diameter (in.) for equivalent cooling rate at: hardenability curve that intersects Three-quarter-radius Half-radius Center the specified hardness at the indicated distance from the quenched end of the Oil at Water at Oil at Water at Oil at Water at hardenability 200 sfm; 200 sfm; 200 sfm; 200 sfm; 200 sfm; 200 sfm; specimen H = 0.5 H = 1.5 H = 0.5 H = 1.5 H = 0.5 H = 1.5

30 HRC 21=2

8617, 4118, 4620, 5120, 1038, 1522, 4419

0.4

1.5

...

1.1

...

0.8

3

4812, 4027, 1042, 1045, 1146, 1050, 1524, 1526, 4028, 6118

0.6

1.8

...

1.2

0.3

0.95

31=2

4720, 6120, 8620, 4032

0.7

2.05

0.5

1.4

0.45

1.1

4

4815, 8720, 4621, 8622, 1050(b)

0.9

2.35

0.7

1.5

0.6

1.3

41=2

46B12, 4817, 4320, 8625, 5046

1.05

2.6

0.8

1.6

0.7

1.45

5

4037, 1541, 4718, 8822

1.2

2.9

0.9

1.8

0.85

1.6

5 =2

94B15, 8627, 4042, 1541, 15B35

1.4

3.2

1.1

1.9

1.0

1.7

6

94B17

...

...

...

...

...

...

6 =2

4820, 1330, 4130, 8630, 1141

1.7

3.8

1.4

2.2

1.25

2.0

7

9130, 5130, 5132, 4047

1.85

...

1.5

2.4

1.35

2.1

71=2

1335, 50B46, 15B37

2.0

...

1.7

2.5

1.5

2.2

8

5135

2.1

...

1.8

2.7

1.6

2.35

91=2

1340

2.5

...

2.2

3.3

1.9

2.7

10

8635, 5140, 4053, 50B40

2.6

...

2.3

3.4

2.0

2.8

11

4640

2.8

...

2.4

3.7

2.15

3.0

1

1


Page 775



12

8637, 1345, 50B44, 5145, 94B30

14

50B50

16

4135, 5147, 8645, 8740

20 22

3.05

...

01 Sep 2005

2.6

3.9

2.3

3.2

...

...

...

...

...

...

3.85

...

3.3

...

2.8

3.85

4063

...

...

3.6

...

...

...

4068, 50B60, 5155, 86B30, 9260

...

...

3.7

...

...

...

24

4137, 5160, 6150, 81B45, 51B60, 8650

...

...

3.85

...

...

...

32

4140

...

...

...

...

...

...

35 HRC 11=2

8617

...

0.9

...

0.8

...

0.45

2

4812, 4118, 4620, 5120, 1038, 1522, 4419, 6118

...

1.2

...

0.9

...

0.65

21=2

4028, 4720, 8620, 4027, 1042, 1045, 1146, 1050, 1524, 1526

0.4

1.5

...

1.1

...

0.8

3

9310, 46B12, 4320, 6120, 8720, 4621, 8622, 8625, 4032, 4815

0.6

1.8

...

1.2

0.3

0.95

31=2

4815, 4817, 94B17, 5046, 1050(b), 4781, 8822

0.7

2.05

0.5

1.4

0.45

1.1

4

8627, 4037

0.9

2.35

0.7

1.5

0.6

1.3

41=2

94B15, 4042, 1541

1.05

2.6

0.8

1.6

0.7

1.45

5

4820, 1330, 4130, 5130, 8630, 5132, 1141, 50B46, 4047, 15B35, 94B17

1.2

2.9

0.9

1.8

0.85

1.6

51=2

1335

1.4

3.2

1.1

1.9

1.0

1.7

6

5135

1.55

3.5

1.2

2.1

1.1

1.85

61=2

15B37

...

...

...

...

...

...

7

8635, 1340, 5140, 4053

1.85

...

1.5

2.4

1.35

2.1

8

4063, 1345, 5145

2.1

...

1.8

2.7

1.6

2.35

81=2

8637

2.2

...

1.9

2.9

1.7

2.45

9

4640, 4068, 50B40

2.35

...

2.0

3.1

1.8

2.6

9 =2

8640, 50B44, 5150

2.5

...

2.2

3.3

1.9

2.7

10

8740, 9260

...

...

...

...

...

...

101=2

4135,50B50

2.7

...

2.35

3.5

2.1

2.9

13

4137

3.25

...

2.8

...

2.45

3.4

16

4140, 6150, 81B45, 86B30

3.85

...

3.3

...

2.8

3.85

5120, 6120

...

0.65

...

0.6

...

0.3

1 =2

4118, 4620, 4320, 4720, 8620, 8720, 1038, 1522, 1526, 4621

...

0.9

...

0.8

...

0.45

2

8622, 8625, 4027, 1045, 1524, 4028, 4718

...

1.2

...

0.9

...

0.65

21=4

1146

0.3

1.3

...

1.0

...

0.7

21=2

4820, 8627, 4032, 1042, 1050

0.4

1.5

...

1.1

...

0.8

3

4037, 8822

0.6

1.8

...

1.2

0.3

0.95

3 =2

4130, 5130, 8630, 5046, 1050(b), 1541

0.7

2.05

0.5

1.4

0.45

1.1

4

1330, 5132, 4042

0.9

2.35

0.7

1.5

0.6

1.3

1

40 HRC 1 1

1


Page 776



01 Sep 2005

41=2

5135, 1141, 4047

1.05

2.6

0.8

1.6

0.7

1.45

5

1335, 50B46, 15B35

1.2

2.9

0.9

1.8

0.85

1.6

1

5 =2

8635, 5140, 4053, 15B37

1.4

3.2

1.1

1.9

1.0

1.7

6

1340, 9260, 4063

1.55

3.5

1.2

2.1

1.1

1.85

61=2

8637, 5145, 1345

1.7

3.8

1.4

2.2

1.25

2.0

7

4640, 4068

1.85

...

1.5

2.4

1.35

2.1

71=2

8640, 5150

2.0

...

1.7

2.5

1.5

2.2

8

4135, 8740, 50B40

2.1

...

1.8

2.7

1.6

2.35

81=2

6145, 9261, 50B44, 5155

2.2

...

1.9

2.9

1.7

2.45

9

4137, 8642, 5147, 50B50, 94B30

2.35

...

2.0

3.1

1.8

2.6

8742, 8645, 5160, 9262

2.5

...

2.2

3.3

1.9

2.7

10 =2

6150, 50B60

2.7

...

2.35

3.5

2.1

2.9

11

4140

2.8

...

2.4

3.7

2.15

3.0

111=2

81B45, 8650, 5152

2.9

...

2.5

3.8

2.25

3.1

12

86B30

...

...

...

...

...

...

13

51B60

3.25

...

2.8

...

2.45

3.4

14

8655

3.45

...

2.95

...

2.6

3.55

15

4142

3.65

...

3.1

...

2.7

3.7

15 =2

8750

3.75

...

3.2

...

2.75

3.8

18

4145, 8653, 8660

...

...

3.45

...

...

...

19

9840, 86B45

...

...

3.45

...

...

...

20

4147

...

...

3.6

...

...

...

24

4337, 4150

...

...

3.85

...

...

...

32

4340

...

...

...

...

...

...

36+

E4340, 9850

...

...

...

...

...

...

4027, 4028, 8625

...

...

...

...

...

...

91=2 1

1

45 HRC 1 1

1 =2

8627, 1038

...

0.9

...

0.8

...

0.45

2

4032, 1042, 1146, 1045

...

1.2

...

0.9

...

0.65

21=2

4130, 5130, 8630, 4037, 1050, 5132

0.4

1.5

...

1.1

...

0.8

3

1330, 5046, 1541

0.6

1.8

...

1.2

0.3

0.95

31=4

1050(b)

0.65

1.9

...

1.3

0.4

1.05

3 =2

1335, 5135, 4042, 4047

0.7

2.05

0.5

1.4

0.45

1.1

4

8635, 1141

0.9

2.35

0.7

1.5

0.6

1.3

5

8637, 1340, 5140, 50B46, 4053, 9260, 15B37

1.2

2.9

0.9

1.8

0.85

1.6

51=2

5145, 4063

1.4

3.2

1.1

1.9

1.0

1.7

6

4135, 4640, 4068, 1345

1.55

3.5

1.2

2.1

1.1

1.85

61=2

8640, 8740, 5150, 94B30

1.7

3.8

1.4

2.2

1.25

2.0

7

4137, 8642, 6145, 9261, 50B40

1.85

...

1.5

2.4

1.35

2.1

1

71=2

8742, 50B44, 5155

2.0

...

1.7

2.5

1.5

2.2

8

8645, 5147

2.1

...

1.8

2.7

1.6

2.35

81=2

4140, 6150, 5160, 9262, 50B50

2.2

...

1.9

2.9

1.7

2.45

9

50B60

2.35

...

2.0

3.1

1.8

2.6

1

9 =2

81B45, 8650, 86B30

2.5

...

2.2

3.3

1.9

2.7

10

5152

2.6

...

2.3

3.4

2.0

2.8


Page 777


11


01 Sep 2005

51B60, 8655

2.8

...

2.4

3.7

2.15

3.0

11 =2

4142

2.9

...

2.5

3.8

2.25

3.1

12

8750

3.05

...

2.6

3.9

2.3

3.2

13

8653, 8660

3.25

...

2.8

...

2.45

3.4

14

9840, 4145

3.45

...

2.95

...

2.6

3.55

16

86B45, 4147

3.85

...

3.3

...

2.8

3.85

17

4337

...

...

3.35

...

...

...

18

4150

...

...

3.45

...

...

...

22

4340

...

...

3.7

...

...

...

26

4161

...

...

...

...

...

...

30

E4340

...

...

...

...

...

...

36

9850

...

...

...

...

...

...

4032, 5132, 1038

...

0.65

...

0.6

...

0.3

1 =2

1335, 5135, 8635, 4037, 1042, 1146, 1045

...

0.9

...

0.8

...

0.45

2

4135, 1541, 15B35, 15B37

...

1.2

...

0.9

...

0.65

1

50 HRC 1 1

1

2 =4

1050(b)

0.3

1.3

...

1.0

...

0.7

21=2

4042

0.4

1.5

...

1.1

...

0.8

3

8637, 5140, 5046, 4047

0.6

1.8

...

1.2

0.3

0.95

31=2

4137, 1141, 1340

0.7

2.05

0.5

1.4

0.45

1.1

4

4640, 5145, 50B46

0.9

2.35

0.7

1.5

0.6

1.3

41=2

8640, 8740, 4053, 9260

1.05

2.6

0.8

1.6

0.7

1.45

5

8642, 4063, 1345, 50B40

1.2

2.9

0.9

1.8

0.85

1.6

51=2

8742, 6145, 5150, 4068

1.4

3.2

1.1

1.9

1.0

1.7

6

4140, 8645

1.55

3.5

1.2

2.1

1.1

1.85

1

6 =2

9261, 50B44, 5155

1.7

3.8

1.4

2.2

1.25

2.0

7

5147, 6150

1.85

...

1.5

2.4

1.35

2.1

71=2

5160, 9262, 50B50

2.0

...

1.7

2.5

1.5

2.2

8

4142, 81B45, 8650

2.1

...

1.8

2.7

1.6

2.35

81=2

5152, 50B60

2.2

...

1.9

2.9

1.7

2.45

91=2

4337, 8750, 8655

2.5

...

2.2

3.3

1.9

2.7

10

4145, 51B60

2.6

...

2.3

3.4

2.0

2.8

101=2

9840

2.7

...

2.35

3.5

2.1

2.9

11

8653, 8660

2.8

...

2.4

3.7

2.15

3.0

11 =2

8645

2.9

...

2.5

3.8

2.25

3.1

12

86B45

...

...

...

...

...

...

13

4340, 4147

3.25

...

2.8

...

2.45

3.4

14

4150

3.45

...

2.95

...

2.6

3.55

20

E4340

...

...

3.6

...

...

...

22

9850, 4161

...

...

3.7

...

...

...

1

1141, 1042, 4042, 4142, 1045, 1146, 1050(b), 8642

...

0.65

...

0.6

...

0.3

11=2

50B46

...

0.9

...

0.8

...

0.45

2

8742, 5046, 4047, 5145

...

1.2

...

0.9

...

0.65

21=2

6145

0.4

1.5

...

1.1

...

0.8

3

4145, 8645, 1345

0.6

1.8

...

1.2

0.3

0.95

1

55 HRC

1

3 =2

86B45, 5147, 4053, 9260

0.7

2.05

0.5

1.4

0.45

1.1

41=2

5150, 4063

1.05

2.6

0.8

1.6

0.7

1.45


Page 778


5


01 Sep 2005

81B45, 6150, 9261, 5155

1.2

2.9

0.9

1.8

0.85

1.6

5 =2

8650, 5152, 4068

1.4

3.2

1.1

1.9

1.0

1.7

6

50B50

1.55

3.5

1.2

2.1

1.1

1.85

61=2

5160, 9262

1.7

3.8

1.4

2.2

1.25

2.0

7

4147, 8750, 8655

1.85

...

1.5

2.4

1.35

2.1

71=2

50B60

2.0

...

1.7

2.5

1.5

2.2

9

8653, 51B60, 8660

2.35

...

2.0

3.1

1.8

2.6

91=2

4150

2.5

...

2.2

3.3

1.9

2.7

17 9850 ... ... 3.35 (a) If based on equivalent hardness, actual bar diameter will be less. (b) High residual alloy

...

...

...

1

The use of Fig. 18 and Table 2 is described in the following example. This method substantially reduces the amount of chart hopping that has in the past been needed to examine all the available steels for the purpose of selecting one. Example 1: Selection of a Steel with 38 mm (1 1=2 in.) Diam Section Equivalent Having 45 HRC at Half-Radius. This example traces the steps needed to select a steel that will harden to 45 HRC at half-radius in a part having a significant section equivalent to a 38 mm (11=2in.) diam bar. First, it is assumed that, to prevent distortion, the quench will be in oil at 60 m/min (200 sfm) (H = 0.5) and that a nonscaling atmosphere will be used for heating to the austenitizing temperature. Therefore, the chart for half-radius in Fig. 18 (c) is applicable. The following steps will then lead to the selection of a steel. First, trace horizontally at the level of 11=2in. diameter to the curve for oil quench at 60 m/min (200 sfm) (curve 5). From the point of intersection with this curve, trace vertically to the x-axis to determine the location on the end-quenched bar that has the same cooling rate as the point at half-radius in the 11=2in. round. This location is 6:5=16from the quenched end of the bar. Then turn to the section of Table 2 that gives the location of 45 HRC on the end-quenched bar for the various H-steels. Here it is found that four steels will produce 45 HRC at 6:5=16from the end of the bar: 8640, 8740, 5150, and 94B30. If some additional hardenability is not undesirable, steels that will produce 45 HRC at 7=16can be included⎯4137, 8642, 6145, and 50B40. Steel 9261 is also in the same category, but it would not be applicable, because it is a spring steel used only when the asquenched hardness must be as high as 50 to 55 HRC. Therefore, eight steels are available that will meet the hardenability requirements of the stipulated specification. From knowledge of other characteristics of these steels, including machinability, forgeability, crackability, distortion, availability, and cost, the selector can decide which of these eight will be the most desirable for the part in question. Scaled Rounds. When values for scaled round bars are desired, Fig. 18 (b), 18 (d), and 18 (f) can be used. However, prediction of results for sizes less than 25 mm (1 in.) in diameter should not be the basis for important decisions involving costly purchases without further checking, because values for these sizes were obtained by extrapolation. Figure 20 shows another correlation for rounds based on the equivalent hardness criterion. In Fig. 20 , cooling conditions from the surface to the center of rounds of various sizes quenched in media ranging in quenching severity from 0.20 to infinity are correlated with Jeh; they are given in 1=16in. units producing the same hardness on the end-quench bar. Figure 20 is especially useful for estimating through-section strength, because the entire hardness profile of the prospective steel (and, to a degree, microstructure as well) can be predicted for rounds with different diameters from one set of end-quench data. Instructions for the procedure are given in the caption. Fig. 20 Correlation of Jeh equivalent hardness positions in end-quenched hardenability specimen and various locations in round bars quenched in oil, water, and brine. The dashed line shows the various positions in 1=2 to 4 in. diam rounds that are equivalent to the 8=16in. distance on the end-quench bar. To determine cross-sectional hardnesses from results of end-quench tests, pick out the end-quench hardness at an appropriate point on the bottom line and extend an imaginary line upward to the curved line that corresponds to the quenching severity needed to obtain that hardness for the given diameter of round.


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Rectangular or Hexagonal Bars and Plate. Except in critical or borderline applications, size relationships for rounds can be applied without correction to square or hexagonal sections. Figures 18 , 19 , and 20 can also be used for rectangular bars in which the ratio of width to thickness (W/T) is less than 4, but the value 1.4 times the thickness should be used as the equivalent round. Large plates cool considerably more slowly than bars. The cooling rate relationships shown in Fig. 21 and 22 apply to these shapes.


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Fig. 21 Correlation of equivalent cooling rates in the end-quench specimen and quenched plates

Fig. 22 Correlation between J ec and center cooling rates in plates quenched at various severities

Tubular Parts. The application of end-quenched hardenability data to the selection of steel for hollow cylindrical sections is based largely on production experience with similar parts. There has been some progress in equating tubular sections to round bars and in developing dimensionless temperature-time charts for long hollow cylinders. Hollomon and Zener (Ref 7) determined by calculation the diameter of solid steel cylinders that, when quenched in a given medium, could be expected to have the same hardness at the center as the minimum hardness in the wall of hollow cylinders when quenched in the same medium. The rule of thumb of doubling the tube wall thickness to obtain the diameter of an equivalent solid bar is a useful first approximation. Estimating Hardenability. When actual end-quenched hardenability data are unavailable, the hardening performance of a steel of given chemical composition can be estimated from calculated hardenability data. The various methods proposed for calculating hardenability are given in the section "Calculation of Hardenability" in this article. Details can be found in Ref 3 and 8, 9, 10, 11, 12, 13.

Use of the Charts The true measure of applicability of any steel to a part requiring heat treatment is the relation of its hardenability to the critical cross section of the part at the time it is heat treated. The term critical cross section refers to that section of the part where service stresses are highest and therefore where the highest mechanical properties are required. For example, if the part is a rough forging 64 mm (21=2in.) in diameter at the critical cross section, which is later machined to 50 mm (2 in.) in diameter, and the finished part must be hardened to three-quarter-radius (that is, 6.4 mm, or 1=4in., deep), then the hardenability of the steel must be such


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01 Sep 2005

that the rough forging will harden 13 mm (1=2in.) deep. Figure 23 shows the correlation between cooling rates along the end-quench hardenability specimen and at four locations in round bars up to 102 mm (4 in.) in diameter for both oil and water quenching at 60 m/min (200 sfm). The curves in Fig. 23 provide data that can be used directly in steel selection. Following is an example of their practical application to a specific problem of steel selection. Fig. 23 Equivalent cooling rates for round bars quenched in water (a) and oil (b). Correlation of equivalent cooling rates in the end-quenched hardenability specimen and quenched round bars free from scale. Data for surface hardness are for mild agitation; other data are for 60 m/min (200 sfm).

Example 2: Use of Hardenability Charts to Verify that 4140H Steel Will Fulfill Hardness Specifications for a 44.45 mm (1.75 in.) Diam Shaft. A shaft 44.45 mm (1.75 in.) in diameter and 1.1 m (31=2ft) long is required in a machine. The engineering analysis indicates that the torsion requirements will approach a maximum of 170 MPa (25 ksi) and that the bending stresses will reach a maximum of 550 MPa (80 ksi). Because several other parts in production in the same plant are being made from 4140H steel, it is desired to know whether 4140H has enough hardenability for this shaft. Because the shear stress in torsion is about one-half that in bending, the latter will be the primary consideration. In bending, stresses approach zero in the neutral axis; therefore, the steel need not be hardened completely to the center. This is helpful


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because the distribution of stress in quenching will decrease the danger of quench cracking and, after tempering, should leave the exterior portion of the shaft in compression. In order to withstand a fatigue load of 550 MPa (80 ksi) in bending, a minimum hardness of 35 HRC is required. For this example, it will be assumed that 35 HRC should be obtained by tempering a structure that, as-quenched, contains at least 80% martensite. From experience with similar parts, it is known that the 80% martensite structure should be present down to the three-quarter-radius position in the shaft. Because 4140H has a minimum carbon content of 0.37%, the first operation on the charts (Fig. 24 ) is to find the as-quenched hardness that corresponds to 0.37% C in an 80% martensite structure. As shown in the top chart of Fig. 24 ⎯the same data as in Fig. 1 (d)⎯this as-quenched hardness is 45 HRC. Fig. 24 Illustration of the use of hardenability data in steel selection

The original question (whether 4140H is appropriate for this part) can now be rephrased to read: Will 4140H provide the required minimum as-quenched hardness of 45 HRC at three-quarter-radius in the 44.45 mm (1.75 in.) diam shaft? To determine the answer to this question, enter the middle chart of Fig. 24 (this is the same as Fig. 23 b) at the diameter level of 44.45 mm (1.75 in.) and more horizontally to an intersection with the 3=4-radius curve. This intersection occurs at the 6:5=16position on the specimen. Then, move down vertically into the bottom chart to an intersection with the curve for minimum hardenability of 4140H. The intersection occurs at 49 HRC. Because no more than 45 HRC is required, 4140H has more than enough hardenability for this part.

H-Steels Hardenability bands are Jominy curves, based on much historical data, that describe the expected hardenability of many grades of carbon and alloy steels. The H-steels are guaranteed by the supplier to meet these limits for specific ranges of chemical composition. In general, the allowable composition ranges of H-steels are slightly wider than those of steels melted to composition specifications in order to accommodate differences in the residual levels and practices of different mills. These steels are designated by the letter "H" following the composition code or preceding the UNS designation. The charts in the following article, "Hardenability Curves," in this Volume show composition limits and hardenability bands


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for SAE-AISI steels that can be purchased on a hardenability basis. The minimums of these bands are summarized in Table 2 . Current steelmaking technologies⎯ladle refining, for example⎯have permitted excellent control of composition. As a result, hardenability specifications much closer than those indicated by the H-bands can often be worked out between purchaser and supplier. The availability of fractional band hardenability is a function of both the base steel composition and the commonality of the grade. When a H-steel is specified, the steel producer shows on the shipping papers or by some other acceptable means the hardenability characteristics of the heat involved. The heat hardenability is shown either by hardness values at specified reference points or at the following distances from the quenched end of the test specimen: 1=16, 1=8, 1=4, 1=2, 3=4, 1, 11=4, 11=2, 13=4, and 2 in. No reading below 20 HRC is reported. The heat hardenability is determined from either a cast or a forged end-quench test bar.

Use of Hardenability Limits H-band limits are shown graphically in the article "Hardenability Curves" in this Volume for convenience in estimating the hardness values obtainable at various locations on the end-quench test bar and for quick comparison of the various H-grades. However, the graphs are not used for specification purposes. Tables appearing with the graphs in the article "Hardenability Curves" in this Volume show the minimum and maximum Rockwell C hardness values at the corresponding distances from the quenched end of the standard end-quenched test specimen for all H-steels. When desirable, the maximum and minimum limits at the 1=16-in. position can be specified in addition to the other two points. When it is necessary to specify more than two points on the hardenability band (exclusive of the maximum and minimum limits at 1=16in.), a tolerance of two points Rockwell C over a 3=16in. portion of either curve (except at 1=16in.) is permitted. Maximum Hardenability Limits As pointed out in the preceding paragraphs, maximum hardenability can sometimes be specified as well as minimum. Although minimum hardenability is significant in relation to the maximum section to be hardened, the maximum hardenability is related chiefly to minimum sections and their tendencies to distort or crack, especially when made from higher-carbon steels. For example, assume that there is a part for which 4137 will not quite provide the necessary minimum hardenability. By changing to 4142, the minimum hardenability will be increased by 4=16at 45 HRC, a worthwhile increase. The danger is that, by increasing carbon to provide greater hardenability, the maximum as-quenched hardness is also increased. Many parts have a thin section sensitive to both maximum hardenability and high as-quenched hardness, and because of this combined effect of higher carbon, such sections often break during quenching. The higher-carbon steels transform to martensite at a lower temperature; that is, the Ms temperature is lower. At this temperature, the steel is less plastic and therefore less able to withstand the strains set up by the volume increase (about 11=2%) when austenite transforms to martensite. Also, the higher-carbon martensites are harder and more brittle and cannot withstand the severe strains set up in quenching as well; therefore, pieces with an unfavorable configuration, such as a shaft with a flange, may develop quench cracks.

Steels for Case Hardening Many alloy steels for case hardening are now specified on the basis of core hardenability. Although the same considerations generally apply as for the selection of uncarburized grades, there are some peculiarities in carburizing applications. First, in a case-hardened steel, the hardenability of both case and core must be considered. Because of the difference in carbon content, case and core have quite different hardenabilities, and this difference is much greater for some steels than for others. Moreover, the two regions have different functions to perform in service. Until the introduction of lean alloy steels such as the 86xx series, with and without boron, there was little need to be concerned about case hardenability, because the alloy content combined with the high carbon content always provided adequate hardenability. This is still generally true when the steels are direct quenched from carburizing, so that the carbon and alloying elements are in solution in the case austenite. In parts that are reheated for hardening and in heavy-sectioned parts, however, both case and core hardenability requirements should be carefully evaluated. The hardenability of the steels as purchased will be the core hardenability. Because these low-carbon steels, as a class, are shallow hardening and because of the wide variation in the section sizes of case-hardened parts, the hardenability of the steel must be related to some critical section of the part, for example, the pitch line or the root of a gear tooth. This is best accomplished by making a part of a steel of known hardenability, heat treating it, and then, by means of equivalence of hardness, relating the hardenability in the critical section or sections to the proper positions on the end-quench hardenability specimens, both base carbon and carburized. Finally, notice that the relationship between the thermal gradient and the carbon (hardenability) gradient during quenching of a carburized part can make a measurable difference in the case depth as measured by hardness. That is, an increase in base hardenability can produce a higher proportion of martensite for a given carbon level, yielding a deeper measured case depth. Therefore, a shallower carbon profile and shorter carburizing time could be used to attain the desired result in a properly chosen steel. Core Hardness. A common mistake is to specify too narrow a range of core hardness. When the final quench is from a temperature high enough to allow the development of full core hardness, the hardness variation at any location will be that of the


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hardenability band of the steel at the corresponding position on the end-quenched hardenability specimen. One way to alter this state of affairs is to use higher-alloy steels. In the commonly used alloy steels having a maximum of 2% total alloy content, the range for the core hardness of sections such as gear teeth is 12 to 15 points HRC. Higher-alloy steels exhibit a narrower range; for example, in 4815 the range is 10 points, while in 3310 it is 8 points. Such steels are justified only for severe service or special applications. In standard steels purchased to chemical composition requirements rather than to hardenability, the range can be 20 points HRC or more; for example, 8620 may vary from 20 to 45 HRC at the 4=16in. position. This 25-point range emphasizes the advantage of purchasing to hardenability specifications to avoid the intolerable variation possible within the ranges for standard chemistry steels. Without resorting to the use of high-alloy steels, another way to control core hardness within narrow limits is to use a final quench from a lower temperature, so that full hardness will be developed in the case without the disadvantage of excessive core hardness. Case-Hardened Steel Applications. In addition to the complexities already mentioned, there are highly variable conditions in heat treating and sometimes differences of opinion, even among qualified engineers. The subject can be simplified to some extent by dividing it into applications involving, first, gears and similar parts and, second, all others. Gears are almost always oil quenched because distortion must be held to the lowest possible level. This means that alloy steels are usually selected⎯which particular alloy is much debated. The lower-alloy steels such as 4023, 5120, 4118, 8620, and 4620, with a carbon range between 0.15 and 0.25%, are widely used and generally satisfactory. The first choice usually would be made from the last two steels mentioned, either of which should be safe for all ordinary applications. The final choice, based on service experience or dynamometer testing, should be the least expensive steel that will do the job. To this list should be added 1524, which, although not classified commercially as an alloy steel, has sufficient manganese to make it oil hardening up to an end-quench correlation point of 3=16. For heavy-duty applications, higher-alloy grades such as 4320, 4817, and 9310 are justifiable if based on actual performance tests. The life testing of gears in the same mountings used in service, to prove both the design and the steel selection, is particularly important. The carbonitriding process extends the use of carbon steels such as 1016, 1018, 1019, and 1022 into the field of light-duty gearing by permitting the use of oil quenching in teeth of eight diametral pitch and finer. Steels selected for such applications should be specified silicon-killed fine-grain in order to ensure uniform case hardness and dimensional control. The core of such gears will, of course, have the properties of low-carbon steel, oil quenched. In the thin sections of fine-pitch teeth, this may be up to 25 HRC. The carbonitriding process is usually limited, for economic reasons, to maximum case depths of approximately 0.6 mm (0.025 in.). Nongear Applications. In other applications, when distortion is not a major factor, the carbon steels described above, water quenched, can be used up to a 50 mm (2 in.) diameter. In larger sizes, low-alloy steels, water quenched, such as 5120, 4023, and 6120, can be used, but possible distortion and quench cracking must be guarded against.

Steel Castings Cast steels have about the same hardenability values as their wrought counterparts of the same composition. Because of their coarser grain, steels in the as-cast condition frequently show higher hardenability than the same steels in the wrought condition, but after they are normalized, hardenability will be more nearly equal. Variation in the hardenability of cast steels is caused by the same factors as in wrought steels. Results from tests on standard end-quench specimens taken from various locations within a heavy casting reveal no significant effect of location on the hardenability of a modified 4032 steel (Fig. 25 ). Additional data on the hardenability of cast steels are given in SAE J434 and J435 and in Ref 14. Fig. 25 Range of results for hardenability specimens cut from various locations in a 4032 alloy modified 1.35% Mn steel casting. Dimensions in illustration given in inches


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REFERENCES 1. C.F. Jatczak, Effect of Microstructure and Cooling Rate on Secondary Hardening of Cr-Mo-V Steels, Trans. ASM, Vol 58, 1965, p 195 2. R.A. Grange, Estimating the Hardenability of Carbon Steels, Metall. Trans., Vol 4, Oct 1973, p 2231 3. M.A. Grossmann, Hardenability Calculated from Chemical Composition, Trans. AIME, Vol 150, 1942, p 227 4. D.V. Doane and J.S. Kirkaldy, Ed., Hardenability Concepts With Applications to Steel, The Metallurgical Society, 1978 5. C.S. Siebert, D.V. Doane, and D.H. Breen, The Hardenability of Steels, American Society for Metals, 1977 6. D.J. Carney, Another Look at Quenchants, Cooling Rates and Hardenability, Trans. ASM, Vol 46, 1954, p 882 7. J.H. Hollomon and C. Zener, Quenching and Hardenability of Hollow Cylinders, Trans. ASM, Vol 33, 1944, p 1 8. J.M. Hodge and M.A. Orehoski, Relationship Between Hardenability and Percentage of Martensite in Some Low-Alloy Steels, Trans. AIME, Vol 167, 1946, p 627 9. I.R. Kramer, S. Siegel, and J.G. Brooks, Factors for the Calculation of Hardenability, Trans. AIME, Vol 167, 1946, p 670 10. A.F. de Retana and D.V. Doane, Hardenability of Carburizing Steels, Met. Prog., Vol 100 (No. 3), Sept 1971, p 65 11. C.F. Jatczak, Trans. AIME, Vol 4, 1973 12. E. Just, New Formulas for Calculating Hardenability Curves, Met. Prog., Vol 96 (No. 5), Nov 1969, p 87 13. J.S. Kirkaldy, Metall. Trans., Vol 4, Oct 1973 14. Materials, Vol 1, 1989 SAE Handbook, Society of Automotive Engineers, 1989 SELECTED REFERENCES • • • • •

A.L. Boegehold, Hardenability Control for Alloy Steel Parts, Met. Prog., Vol 53 (No. 5), May 1948, p 697 D.V. Doane and J.S. Kirkaldy, Ed., Hardenability Concepts With Applications to Steel, The Metallurgical Society, 1978 C.F. Jatczak, Determining Hardenability From Composition, Met. Prog., Vol 100 (No. 3), Sept 1971, p 60 J.L. Lamont, Iron Age, 14 Oct 1943 R.A. Rege, P.E. Hamill, and J.M. Hodge, The Effects of Geometry on the Cooling of Plates and Bars, Trans. ASM, Vol 62, 1969, p 333 • Report OSRD 3743, National Defense Research Committee • J.M. Tartaglia and G.T. Eldis, Core Hardenability Calculations for Carburizing Steels, Metall. Trans. A, Vol 15A (No. 6), June 1984, p 1173−1183 • E.W. Weinman, R.F. Thomson, and A.L. Boegehold, Correlation of End-Quenched Test Bars and Rounds, Trans. ASM, Vol 44, 1952, p 803


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Hardenability Curves

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ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Hardenability of Carbon and Low-Alloy Steels Hardenability Curves HARDENABILITY CURVES for more than 80 types of carbon and alloy H-band steels comprise this article (Ref 1, 2). The tabular data used to compile these curves are also included with each graph. Values from these tables are used for specification purpose, and SAE recommends choosing two points to designate the hardenability.

Hardenability Curves The two points may be designated in any one of the ways listed below and illustrated in Fig. 1 : • The minimum and maximum hardness values at any desired distance. This method is illustrated as points A-A and is specified as J43 to J54 = 3=16in. Obviously, the distance selected would be that distance on the end-quench specimen that corresponds to the section used by the consumer • The minimum and maximum distances at which any desired hardness value occurs. This method is illustrated as points B-B and would be specified as J39 = 4=16to 9=16in. • Two maximum hardness values at two desired distances, illustrated as points C-C and specified as J50 = 5=16in. (max), J34 = 12 =16in. (max) • Two minimum hardness values at two desired distances, illustrated as points D-D and specified as J35 = 5=16in. (min), J21 = 16 =16in. (min) • Any minimum hardness plus any maximum hardness, E-E, specified as J37 max = 10=16, J32 min =6=16 It should be noted that each H-band hardenability limit curve is presented graphically and in tabular form, in both metric and English units. Fig. 1 Typical hardenability curve shown with English units. See text for discussion of designated points.


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REFERENCES 1. Materials, Vol 1, 1989 SAE Handbook, Society of Automotive Engineers, 1989 2. "Alloy, Carbon and High Strength Low Alloy Steels: Semifinished for Forging; Hot Rolled Bars and Cold Finished Bars, Hot Rolled Deformed and Plain Concrete Reinforcing Bars," Steel Products Manual, American Iron and Steel Institute, March 1986 SAE/AISI 1038H UNS H10380. Heat-treating temperatures recommended by SAE ÃNormalize (for forged or rolled specimens only): 870 °C (1600 °F) ÃAustenitize: 845 °C (1550 °F)


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SAE/AISI 1045H UNS H10450. Heat-treating temperatures recommended by SAE ÃNormalize (for forged or rolled specimens only): 870 °C (1600 °F) ÃAustenitize: 845 °C (1550 °F)


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SAE/AISI 15B21H UNS H15211. Heat-treating temperatures recommended by SAE ÃNormalize (for forged or rolled specimens only): 925 °C (1700 °F) ÃAustenitize: 925 °C (1700 °F)


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SAE/AISI 4027H UNS H40270. SAE/AISI 4028H UNS H40280 Heat-treating temperatures recommended by SAE ÃNormalize (for forged or rolled specimens only): 900 °C (1650 °F) ÃAustenitize: 870 °C (1600 °F)


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SAE/AISI E4340H UNS H43406. Heat-treating temperatures recommended by SAE ÃNormalize (for forged or rolled specimens only): 870 °C (1600 °F) ÃAustenitize: 845 °C (1550 °F)


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Sheet Formability of Steels

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Fabrication Characteristics of Carbon and Low-Alloy Steels Sheet Formability of Steels Revised by W.G. Granzow, Armco, Inc., Research & Technology STEEL SHEET is widely used for industrial and consumer products, partly because it is relatively strong, easily joined, and readily available at moderate cost. However, it is not these factors, but the formability of steel sheets that is primarily responsible for its wide use. The term formability is commonly used to described the ability of a steel to maintain its structural integrity while being plastically deformed into various shapes. However, not all shapes require the same forming characteristics, and a steel that has excellent formability in one application may exhibit poor formability in another application. In practice, therefore, formability must be optimized by selecting a grade of steel that has the forming characteristics needed to make the required shape. These forming characteristics are normally estimated from an analysis of the mechanical properties of steel, which are determined by uniaxial tensile tests. Although this type of test does not simulate any commercial forming operations, the test results have been universally used for many years to evaluate formability, and some understanding of them is essential to the understanding of sheet steel formability. Sheet metal forming methods are described in detail in Forming and Forging, Volume 14 of ASM Handbook. Examples of formed parts that require different forming characteristics in the steel are shown in Fig. 1 . Part A was formed by drawing; that is, all the metal that was required to form the part from a flat blank came from the flanges. This shape requires that the steel have a high plastic-strain ratio, or r value, which determines the resistance of steel sheet to thinning during forming operations. Part C was formed by stretching; the flange on the blank was clamped during forming, and all of the metal that was required to form the part came from reducing the thickness of the metal. This type of part requires good ductility in the steel. However, the r value should be low, because high r values can cause failures of stretched parts. Part B has failed in plane strain, which is a type of stretching. Parts that develop this strain condition, such as automotive panels, require good ductility. Fig. 1 Parts that required different forming characteristics in the steel sheet

This article discusses the mechanical properties and formability of steel sheet, the use of circle grid analysis to identify the properties of complicated shapes, and various simulative forming tests. It covers the effects of steel composition, steelmaking practices, and metallic coatings, as well as the correlation between microstructure and formability. A guide to the selection of steel sheet is also included.

Mechanical Properties and Formability The mechanical properties of steel sheet that influence its forming characteristics, either directly or indirectly, can be measured by uniaxial tension testing, such as that described in ASTM E 8. The tensile test results of particular interest include the yield strength, ultimate tensile strength, total elongation, uniform elongation, yield point elongation, plastic-strain ratio, planar anisotropy, and the strain-hardening exponent. Uniaxial tensile tests may be made with specimens obtained from longitudinal, diagonal, transverse, or other orientations relative to the rolling direction. Typical mechanical properties for common grades of hot-rolled and cold-rolled steel sheets are given in Tables 1 and 2 . Table 1 Typical mechanical properties of hot-rolled steel sheet Yield


Tensile

Elongation

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strength Type or quality

Special feature

MPa

strength

ksi

MPa

ksi

01 Sep 2005

in 50 mm (2 in.), %

Hardness, HRB

Strain-harde Plastic-strai n ning ratio, rm exponent, n

Commercial

Standard properties

262

38

359

52

30

55

0.15

0.9

Drawing (rimmed)

Improved properties

241

35

345

50

35

50

0.18

1.0

Drawing (special killed)

Nonaging

241

35

345

50

40

50

0.20

1.0

Medium strength

Inclusion shape control

345

50

414

60

25

70

0.15

0.9

High strength

Inclusion shape control

552

80

620

90

15

90

...

...

Table 2 Typical mechanical properties of cold-rolled steel sheet Yield strength Type or quality

Special feature

MPa

ksi

Tensile strength MPa

ksi

Elongation in 50 mm (2 in.), %

Hardness, HRB

Strain-hardeni Plastic-str ain ng ratio, rm exponent, n

Commercial

Standard properties

234

34

317

46

35

45

0.18

1.0

Drawing (rimmed)

Stretchable

207

30

310

45

42

40

0.22

1.2

Drawing (special killed)

Deep drawing

172

25

296

43

42

40

0.22

1.6

Interstitial free

Extra deep drawing

152

22

317

46

42

45

0.24

2.0

Medium strength

Formable

414

60

483

70

25

85

0.20

1.2

High strength

Moderately formable

689

100

724

105

10

25(a)

...

...

(a) HRC

Yield strength of steel sheet is indicative of both formability and strength after forming. Several types of yielding behavior are observed in steel sheet (see Fig. 2 ). When yield point elongation occurs, the lowest value observed during discontinuous yielding is reported as the yield strength. In the absence of an abrupt change in the load-extension curve, the stress at 0.2% offset or 0.5% extension under load is reported as the yield strength. Fig. 2 Load-extension curves for steel sheet having the same yield strength, but different characteristic behavior. (a) Annealed soft-rimmed or aluminum-killed steel; yield strength is the lowest stress measured during yield point elongation. (b) Lightly temper-rolled rimmed steel; stress at the jog in the curve is reported as yield strength. (c) and (d) Temper-rolled low-carbon steel. May be rimmed, aluminum-killed, or interstitial-free steel with no detectable yield point. The yield strength is calculated from the load at 0.2% offset (c) or from the load at 0.5% extension (d). (e) Rimmed steel with a yield point elongation due to aging at room temperature for several months. The yield strength is the lowest stress measured during yield point elongation.

In forming plain carbon steel sheet, a yield strength of 240 MPa (35 ksi) or more increases the likelihood of excessive springback and breakage during forming. However, the use of material have a yield strength of less than 140 MPa (20 ksi) may result in parts with insufficient strength levels. High-strength, formable sheet steels have been developed for applications for which increased strength or reduced weight, in addition to moderate formability, are required. The yield strengths of these steels generally range from 345 to 690 MPa (50 to


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100 ksi). Total Elongation. After fracture, the tensile specimen is pieced together, and the length between gage marks is measured. In this manner, elongation is calculated and reported as a percentage of the original gage length, which is usually 50 mm (2 in.). (A gage length of 200 mm, or 8 in., may be used for heavier-gage metals.) Specimens of sheet metals used for tensile tests usually have short, parallel-sided, reduced sections, but slightly tapered reduced sections are sometimes used to control the location of necking and fracture. Values of elongation resulting from tests of different specimens of the same material may vary because of differences in gage length, sheet thickness, edge preparation and finish, test methods, or other factors. Typical values of the amount of elongation in 50 mm (2 in.) are listed in Tables 1 and 2 for common formable grades of steel sheet. Generally, an elongation of 35 to 45% in 50 mm (2 in.) is normal for conventional low-carbon steels, with higher values indicating better formability. Uniform Elongation. The total elongation of a sheet tensile specimen comprises two parts, uniform elongation and postuniform elongation. For a material that follows the power relationship for hardening (σ = Kεn), the uniform elongation (measured in true strain) is equal to the strain-hardening exponent, n. The postuniform elongation depends on both the strain-hardening behavior and the strain rate sensitivity response of the metal to the applied stress. When a neck forms, the strains and strain rate within the neck are greater than in the outside regions, and increased strain hardening may offset the weakening due to the reduced cross-sectional thickness, causing a shift of deformation to regions outside the neck. The engineering elongation to maximum load, eu, is related to the strain-hardening exponent, n, by the equation: n = ln (1 + eu) (Eq 1) Typical values of eu for low-carbon steels range from 20 to 30%. The eu and associated n values indicate the work-hardening rate of sheet metals and, thus, the capability of the metal to deform in stretch, plain-strain, and bending deformation modes. Other factors, such as strain-rate sensitivity, can enhance or detract from the capability of a metal to be formed into a part. For example, the n values of low-carbon steel and 1100-O grade aluminum are about the same; however, both the total elongation and the forming limit of aluminum are considerably lower than those of low-carbon steel because aluminum has a negative value of m, the strain-rate sensitivity in response to the applied flow stress:

(Eq 2) where σ is the flow stress and ²_is the strain rate. For sheet metals that fail by local necking, uniform elongation may not give a true estimate of formability. Estimates based on total elongation are often considered more reliable. Yield point elongation is the portion of total elongation that occurs during discontinuous yielding at the yield stress. It is accompanied by the formation of surface defects known as Lüders lines, or stretcher strains, which are considered imperfections in many applications of steel sheet because of their unsightly appearance. Yield point elongation during tensile testing indicates that Lüders lines are likely to occur during forming. Yield point elongation requires the presence of interstitial residual alloying elements, particularly carbon or nitrogen; consequently, low-interstitial steels do not exhibit this effect. Yield point elongation can be suppressed by temper rolling the steel sheet at the mill. However, unless the nitrogen has been combined with another element (usually aluminum), the steel will age harden after a period that varies from a few hours to a year or more (depending on storage temperature and other factors). Aged steels can be used in most forming operations, provided they are roller leveled or flex immediately before fabrication, although these methods are less effective than temper rolling. Plastic-strain ratio, r, describes the resistance of steel sheet to thinning during forming operations. This is the ratio of the true strain in the width direction, εw, to the true strain in the thickness direction, εt, of plastically strained sheet metal:

(Eq 3) The plastic-strain ratio is related to the crystallographic orientation of low-carbon steels. A standard method for determining r by using the tension test is given in ASTM E 517. The value will vary with test direction (relative to the coil rolling direction) in anisotropic metals. An average value, rm, (sometimes designated r¹ ), represents the normal plastic anisotropy of the steel sheet:

(Eq 4) Hot-rolled and normalized cold-rolled steels are generally isotropic (rm of 1.0). Rimmed steels usually have an rm of 1.2, but this value may be higher in special cases, as with some low-manganese low-sulfur products. Aluminum-killed steels will be more anisotropic, with rm of 1.6. Higher values (up to 2.5) may be attained by controlling composition and processing. The upper limit


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for commercial steels is about 3.0, although values near 3.0 are seldom achieved. Interstitial-free steels tend to have the highest rm at approximately 2.0. The rm value predicts the ability of metals to deform in draw. Planar anisotropy may be reported as:

(Eq 5) Planar anisotropy is a measure of the amount of high points, or ears, that will develop on the edges of deep-drawn cylindrical cups or similar parts. High points in the rolling and transverse directions are noted when ∆r is positive (for low-carbon, drawing-quality, aluminum-killed steel sheet). For some high-strength low-alloy steels, ∆r is negative, and earing occurs at 45° to the rolling direction. For most applications, values of ∆r near 0 are preferred, because such values imply a minimal tendency to form ears when metals are drawn into cylindrical cups. The strain-hardening exponent, n, is the slope of the true stress-true strain curve, when plotted on logarithmic coordinates. A significant portion of the curve is nearly a straight line for many low-carbon steels. The data are assumed to fit the equation: σ = Kεn (Eq 6) The n value will normally be approximately 0.22 for low-carbon steels used to form complex-shape parts. Higher values (up to 0.26) indicate improved capabilities to deform in stretch. Freshly rolled rimmed steels generally have n values comparable to those of aluminum-killed steels. After aging, values of n for rimmed steels are less than those for aluminum-killed steels. Some low-carbon steels that are not fully processed for formability, especially hot-rolled grades, will have n values as low as 0.10, but most of the formable grades will have n values above 0.14. The effects of different n values on strain distribution in critical regions of a specific formed part are shown in Fig. 3 . Parts formed from steel sheet with a low n value (0.21) may undergo excessive thinning and fracture in critical regions. Identical parts formed from sheet with a higher n value (0.23) frequently will be strong enough in the critical areas to transfer strain to adjacent areas, thereby avoiding failure during forming. Fig. 3 The major strain ε1 in the critical region of a formed part is more uniformly distributed for the steel having the higher value of n. One of these two parts (which are identical except for the n value of the steel selected) was strained to the point of excessive thinning; the other, made from steel with the higher n value, showed no inclination to fracture. Source: Ref 1

Circle Grid Analysis Uniaxial tension tests determine the mechanical properties of steel sheet under closely controlled, frictionless conditions, which are different from the conditions that normally occur during sheet metal forming operations. However, experience has shown that test results often correlate with the ability of a steel to be successfully formed into parts. The relationship is not always a simple one. For example, the parts shown in Fig. 1 require the steel to exhibit different r and n values for successful processing. With practice, the mechanical properties that are required to form simple shapes like those shown can be estimated, but the properties required to form more complicated shapes can best be determined using circle grid analysis. To use circle grid analysis, a pattern consisting of uniformly spaced, uniformly sized circles (usually 2.5 or 5.0 mm, or 1.0 or 0.2 in., in diameter) is etched onto a flat sheet of steel that is to be formed, and the steel is then processed into the desired shape. The circles change into ellipses with the deformation of the gridded steel blank. The ellipses are then measured to determine the maximum dimension (major strain) and the minimum dimension (minor strain). These strains are then plotted on a forming limit diagram (Fig. 4 ). The location and magnitude of the plotted points on the diagram indicate the severity of the forming operation. Their position on the right (stretch) side or left (draw) side of the diagram indicates the strain condition in the steel, where 0%


Page 878



01 Sep 2005

minor strain indicates a plane-strain condition. Experience has shown that breakage associated with critical strains on the right, or stretch, side may be resolved by increasing the total elongation in the steel. Breakage associated with strains on the left, or draw, side may be eliminated by increasing the r-value of the steel. The magnitude of all of the strains can be reduced by changing the sheet metal forming conditions. Fig. 4 A forming limit diagram (FLD). Strains in the critical zone and above it will result in excessive breakage. Strain conditions on the left side require high r values, while strain conditions on the center and right side require good ductility (a high percentage elongation in tensile tests).

The diagram in Fig. 4 is for 0.914 mm (0.036 in.) thick drawing-quality steel sheets. The critical area curves will be lower for thinner steels and higher for thicker steels. The curves can also be adjusted to accommodate commercial quality steels and other grades with lower n values than drawing-quality steels. The adjustments are made by moving the plane-strain intercept, or the point at which the lower curve crosses the 0% minor strain line, according to Fig. 5 . Fig. 5 Relationship between the plane-strain intercept on a forming limit diagram (FLD0) and the strain-hardening exponent as a function of thickness. FLD0 depends only on thickness for values n greater than 0.21. Source: Ref 2

The plotted points on a forming limit diagram show the magnitude of the strains that develop in steel sheets that are processed


Page 879



01 Sep 2005

into a particular shape. Although the magnitude of the strains will be slightly different in steels with different r values, the plotted points generally show the strain conditions that will occur in any steel that has been used successfully to make a particular part. Therefore, in theory, the quality of the steel that is used to perform the analysis should not be important. In practice, it is advantageous to determine the mechanical properties of the steel when the circle grid analysis is made because excessive strains exhibited by formable steel are a clear indication that changes in processing will be required.

Simulative Forming Tests Numerous tests have been developed to observe and compare the performance of steel sheets under actual forming conditions. The best known are the Olsen, Erichsen, and Swift cup tests, and the limiting dome height (LDH) test. The Olsen and Erichsen cup tests measures steel-stretching performance, the Swift cup test measures steel-drawing performance, and the LDH test measures the performance of a steel under plane-strain conditions. The Olsen cup test measures the maximum penetration of a 22 mm (0.875 in.) diam hemispherical punch into a clamped, flat blank of steel. The punch depth at failure is the Olsen cup value. The Erichsen test is similar, but the punch diameter is 20 mm (0.790 in.). Experience has shown that a conscientious operator using a specific testing machine can use the Olsen or Erichsen cup test results to evaluate the ductility of the steel. However, the correlation between the test results and the steel performance in many sheet metal forming operations has not been good. The Swift cup test determines the maximum diameter circular blank that can be successfully drawn into a flat-bottom cup. The die is usually 100 mm (4 in.) in diameter, although other dimensions are also used. The results are expressed as the limiting draw ratio (LDR):

The LDR is a measure of the drawability of the steel, and the test results correlate well with the r value. The LDH test measures the resistance of a steel to failure under plane-strain conditions. The test sample is a 180 mm (7 in.) long rectangle with an experimentally determined width (~135 mm, or 5.25 in.) that is clamped and stretched to failure over a 100 mm (4 in.) diam hemispherical punch. The sample width is the width that is found to produce the minimum height in the test; this minimum height is the LDH value for that material. The LDH value has been shown to correlate well with the performance of steel sheets in stretch-type automotive body panels, which normally fail under plane-strain conditions. However, it does not correlate with steel performance on deep-drawn panels. All simulative forming tests are affected by friction between the sheet metal and the tooling. This creates problems with sample and tooling preparation because the factors that affect friction in sheet metal forming are poorly understood. In practice, the samples and tooling are usually cleaned with a light mineral oil and wiped dry before testing. This treatment provides sufficient lubricity to prevent a "stick-slip" effect, while minimizing the effect of the lubrication.

Effects of Steel Composition on Formability Low-carbon sheet steels are generally preferred forming. These steels typically contain less than 0.10% carbon and less than 1% total intentional and residual alloying elements. The amount of manganese, the principal alloying addition, normally ranges from 0.15 to 0.35%. Controlled amounts of silicon, niobium, titanium, or aluminum may be added either as deoxidizers or to develop certain properties. Residual elements, such as sulfur, chromium, nickel, molybdenum, copper, nitrogen, and phosphorus, are usually limited as much as possible. In steelmaking shops, these amounts are based on the quality of sheet being produced. Alloy sheet steels (including high-strength low-alloy grades) however, contain specified amounts of one or more of these elements. Carbon content is particularly significant in steels that are intended for complex forming applications. An increase in the carbon content of steel increases the strength of the steel and reduces its formability. These effects are caused by the formation of carbide particles in the ferrite matrix and by the resulting small grain size. The amount of carbon in steel sheet is generally limited to 0.10% or less to maximize the formability of the sheet. Manganese enhances the hot-working characteristics of the steel and facilitates the development of the desired grain size. Some manganese is also necessary to neutralize the detrimental effects of sulfur, particularly for hot workability. Typical manganese contents for low-carbon steel sheet range from 0.15 to 0.35%; manganese contents up to 2.0% may be specified in high-strength low-alloy steels. When the sulfur content of the steel is very low, the manganese content also can be low, which allows the steel to be processed to develop high r values. Phosphorus and sulfur are considered undesirable in steel sheet intended for forming, drawing, or bending because their presence increases the likelihood of cracking or splitting. Allowable levels of phosphorus and sulfur depend on the desired quality level. For example, commercial-quality cold-rolled sheet must contain less than 0.035% P and 0.040% S. For more applications, phosphorus may be added to the steel to increase the strength. Sulfur usually appears as manganese sulfide stringers in the microstructure. These stringers can promote splitting, particularly whenever an unrestrained edge is deformed. Silicon content in low-carbon steel varies according to the deoxidation practice employed during production. In rimmed steels (so called because of the rimming action caused by outgassing during solidification from the molten state), the silicon content is generally less than 0.10%. When silicon rather than aluminum is used to kill the rimming action, the silicon content


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01 Sep 2005

may be as high as 0.40%. Silicon may cause silicate inclusions, which increase the likelihood of cracking during bending. Silicon also increases the strength of the steel and thus decreases its formability. Chromium, nickel, molybdenum, vanadium, and other alloying elements are present in low-carbon steel only as residual elements. With proper scrap selection and control of steelmaking operations, these elements are generally held to minimum amounts. Each of these elements increases the strength and decreases the formability of steel sheet. High-strength low-alloy steels may contain specified amounts of one or more of these elements. Copper is generally considered an innocuous residual element in steel sheet. The strengthening effect of copper is almost negligible in typical residual amounts of less than 0.10%. However, copper is added to steel in amounts exceeding 0.20% to improve resistance to atmospheric corrosion. Niobium strengthens high-strength low-alloy steel through the formation of niobium carbides and nitrides. It can also be used either alone or in combination with titanium to develop high r values in interstitial-free steels. These alloying elements remove the interstitial elements carbon and nitrogen from solid solution. Consequently, the steel shows no yield point elongation. Titanium is a strong carbide and nitride former. It helps develop high r values and eliminates yield point elongation and the aging of cold-rolled annealed steel sheet. Titanium streaks may be a problem in some grades, especially in the form of surface defects in exposed applications. Aluminum is added to steel to kill the rimming action and thus produce a very clean steel known as an aluminum-killed, or special-killed, steel. Aluminum combines with both the oxygen and nitrogen to stop the outgassing of the molten steel when it is added to the ladle or mold. Aluminum also aids the development of preferred grain orientations to attain high r values in cold-rolled and annealed steel sheet. Elongated grains of an approximate ASTM 7 size are found in most well-processed aluminum-killed steels. Because the aluminum combines with the nitrogen, the steel is not subject to strain aging. Nitrogen can significantly strengthen low-carbon steel. It also causes strain aging of the steel. The effects of nitrogen can be controlled by deoxidizing the melt with aluminum. Cerium and other rare earth elements may be added to steel to change the shape of manganese sulfide inclusions from being needlelike or ribbonlike to being globular. Globular inclusions reduce the likelihood of cracking if the sheet is formed without restraining the edges. Oxygen content of molten steel determines its solidification characteristics in the ingot. Excessive amounts of oxygen impede nitride formation and thus negate the effects of alloying elements added to minimize strain aging. Deoxidizers such as silicon, aluminum, and titanium will control the oxygen content. When oxygen combines with these deoxidants, complex nonmetallics are formed. Although most nonmetallics dissolve in the slag, some may become trapped in the steel, causing the surface defects of seams and slivers.

Effects of Steelmaking Practices on Formability The formability of steel sheet is determined to a great extent by the steelmaking practices employed in manufacturing. The user of steel sheet normally specifies certain characteristics for the sheet, thus ensuring that the material can be formed in a predictable manner. Adherence to these specifications also implies that the producer of the sheet has observed whatever steelmaking practices are necessary to enable the product to perform as indicated. The user can specify either hot-rolled or cold-rolled sheet, and he must select an appropriate quality designation. Sometimes deoxidation practice is specified. The user has some latitude in choosing the surface finish of the sheet. Of course, the user specifies the dimensions and tolerances of the sheet and the type of edge to be supplied. Hot-rolled steel is rolled to its final thickness in an elevated-temperature process. The finishing temperature is determined by the composition of the steel and the desired properties. In the as-hot-rolled condition, the steel has a dark-gray oxide coating on its surface, which offers limited corrosion protection as long as it is undisturbed. However, the oxide flakes off during forming and may be undesirable around the press. Because the oxide coating also interferes with steel surface lubricants, it should be removed before the final finishing of most formed parts. Hot-rolled steel may be ordered pickled (using either hot sulfuric or hydrochloric acid to remove the oxide) and oiled to inhibit in-transit rusting. Hot-rolled steel in the as-pickled condition will show stretcher strains or Lüders lines on the surface after forming. Whenever surface appearance is important, the steel should be ordered with a temper-rolled surface (skin pass of less than 2% cold reduction) to reduce this tendency. If aging is a problem because of storage requirements, special-killed hot-rolled steel should be ordered. There is no preferred grain orientation providing high r values in hot-rolled steel, but improved grain size and resistance to longitudinal splitting may be attained by closely controlling chemical composition, which differs between commercial- and drawing-quality hot-rolled steel. Higher strength levels (when necessary for the part being formed) are obtained by alloy additions and processing controls to develop improved structure. Because higher strength is associated with forming problems such as lower ductility, increased springback, and longitudinal bend failures, only high-strength low-alloy steels designed for improved formability should be used in structural parts made by press operations. Mechanical properties of several types of hot-rolled sheet are given in Table 1 . Cold-rolled steel sheet forming is produced by the cold reduction of hot-rolled pickled coils, followed by annealing and possibly additional processing, such as temper rolling. Class 1 (E, exposed) should be ordered when a controlled surface finish is required. Class 2 (U, unexposed) is intended for applications in which surface appearance is not of primary importance. Both classes are available as commercial-quality, drawing-quality, or drawing-quality special-killed cold-rolled steel. Mechanical properties of cold-rolled steel sheet are given in Table 2 . Most cold-rolled steels exhibit yield point elongation in the as-annealed condition. This appears as Lüders lines, or stretcher


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strains, on the surface of formed parts (for example, flat areas near the corners of pan-shape draws) that have been subjected to moderate forming operations. The yield point elongation may be removed by temper rolling the annealed coils. Because temper coiling strengthens the steel and reduces its ductility, it is usually limited to 0.5 to 1.5% elongation of the strip. Temper rolling under tension is more effective than flex rolling or roller leveling for eliminating yield point elongation because the steel is more uniformly strained through the thickness. These latter methods are sometimes used the plants of fabricators because the equipment is less expensive and because it permits the use of aged coils of rimmed steels that may show strain on the surface of formed parts. In addition to many as-processed surface finishes, cold-rolled sheet may be ordered with a metallic coating that provides corrosion protection or a decorative finish that reduces the manufacturing costs of parts such as appliances or building panels. Rimmed steels are available as both hot-rolled and cold-rolled products. The rimming action caused by outgassing during solidification produces a relatively pure iron layer on the surface of the ingot. Thus, rimmed steel generally has a better surface finish than killed steel. After the annealing treatments used to regain ductility in the product following cold reduction to final thickness, rimmed steels must be temper rolled to prevent the formation of Lüders lines during forming. The two available quality levels of rimmed steel are achieved by controlling chemical composition and annealing practice. Commercial quality is standard, whereas drawing quality is produced under stricter tolerance levels for impurities and is given a longer anneal to ensure uniformity throughout the coil, as well as good formability. Rimmed steels are more suited to stretch-type deformation than to deep drawing, for which aluminum-killed steels are generally recommended. Rimmed steels will age after a period of time following temper rolling. Consequently, there is a time limit on any performance guarantee on drawing-quality rimmed steel. Aluminum-killed steels are deoxidized with aluminum and, possibly, with silicon. As already mentioned, use of aluminum results in a very clean steel, known as aluminum-killed or drawing-quality special-killed steel. Exceptional resistance to thinning through the sheet thickness (as measured by the plastic strain ratio, r) can be developed through the controlled processing of these steels. Because the pure iron skin characteristic of rimmed steel does not exist in aluminum-killed steel, surface imperfections may occasionally be encountered on aluminum-killed sheet. Both class 1 and class 2 drawing-quality aluminum-killed steels are produced. It should be noted that some aluminum-killed steels that cannot meet the formability requirements for drawing-quality sheet are sold as commercial-quality steel. Interstitial-free steel is vacuum degassed to reduce the amounts of the interstitial elements carbon, nitrogen, and oxygen. It is usually processed to achieve high values of rm (~2.00). This type of steel is not subject to strain aging at any stage of processing or manufacture; it exhibits no yield point elongation. Interstitial-free steel can withstand deeper draws with less breakage than other grades of steel sheet, and coated products made from it generally retain excellent formability. Surface finish may be specified for cold-rolled steel sheet. The need for uniformity among parts that must have matching surface finishes (such as automobile fenders and hoods), even when made from different materials, often dictates the sheet finish. A surface roughness of 0.8 to 1.5 µm (30 to 60 µin.) for average peak height and two to six peaks per millimeter (50 to 150 peaks per inch) is considered standard for cold-rolled steel sheet. The surface finish is determined by the finish applied to the cold-mill rolls and the temper-mill rolls. Roll finishes are obtained by shot blasting or electroetching a ground roll surface so that the roll is roughened sufficiently to transfer the pattern to the sheet. As these rolls are used, their finish tends to become smoother; there may be a consequent change in appearance among coils, and press performance may vary slightly. A rougher sheet surface tends to hold lubricant better and resists galling and cold welding to die surfaces during forming. For parts requiring little forming, a smoother and often preferred finish can be attained when roughness is minimized.

Correlation Between Microstructure and Formability The formability of steel sheet is related to various microstructural features of the sheet. For example, grain size and shape, grain orientation relative to the rolling direction, and the various microconstituents present in the steel are reflected in its forming behavior. Grain size of steel sheet influences formability in two opposing ways. Petch (Ref 3) has shown that the yield strength of low-carbon steel varies inversely with the square root of the grain diameter. Fine-grained steels are quite strong, but they have low strain-hardening exponents and limited formability. Blickwede (Ref 4) has shown (Fig. 6 ) that rm decreases as grain size decreases. Coarse-grained steels have better formability, but the roughened surface (called orange peel) that results from stretching steel with grain sizes that are below ASTM 5 is unacceptable for many applications. Grain sizes of ASTM 7 or 8 are usually a good compromise between formability and surface appearance. High-strength low-alloy steels, however, are usually produced with extremely small grain sizes (as small as ASTM 12) to increase both strength and toughness. Fig. 6 Variation in rm with grain size for four low-carbon sheet steels. Steels were cold reduced 70% and annealed. Source: Ref 4


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Figure 7 shows the effect of abnormally large surface grains on surface appearance after forming. It should be noted that the effect of the large surface grains is visible on the opposite surface of the sheet. Fig. 7 Forming behavior of decarburized rimmed steel sheet (1.12 mm, or 0.044 in., thick) containing normal grain size distribution and abnormally large surface grains, which resulted from a change from normal manufacturing practice. (a) Cross section of test cup made from normally manufactured steel sheet. Grain size ASTM 6 throughout. (b) Cross section of test cup made from steel sheet containing abnormally large surface grains. Grain size ASTM 3 at one surface and ASTM 7 elsewhere. (a) and (b) both 100×, 3% nital etch. (c) Outside surface of test cup made from normally manufactured steel sheet. (d) Outside surface of test cup made from steel sheet having abnormally large grains on outside surface of cup. Note pronounced orange peel effect. (e) Outside surface of test cup made from steel sheet having abnormally large grains on inside surface of cup. (c), (d), and (e) all 2=3×


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Grain shape of the ferrite can also affect sheet formability. Rimmed and hot-rolled aluminum-killed steels generally have equiaxed grains. Cold-rolled aluminum-killed steels, when properly processed, generally exhibit pancake-shape ferrite grains (Fig. 8 ). This grain shape is associated with the preferred grain orientation that is responsible for the excellent formability of aluminum-killed steels. Fig. 8 Low-carbon steel, cold rolled 65%, showing the grain structure in the rolling plane (R), the longitudinal plane (L), and the transverse plane (T). RD, rolling direction

The microconstituents that are found in low-carbon steel at room temperatures include iron carbides and various nonmetallic inclusions. The most common inclusions are sulfides, silicates, and oxides. Aluminum-killed steels will also contain submicroscopic particles of aluminum nitrides. These microconstituents can affect the formability of steel sheet by altering its strength. Alloying elements that dissolve in ferrite strengthen the steel appreciably, thereby reducing its formability. Nonmetallic inclusions may form a distinctive pattern of stringers that reflects the processing history from ingot to sheet. These elongated particles affect the formability of sheet primarily because they encourage cracking at the edge of a part during forming.

Effect of Metallic Coatings on Formability


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Low-carbon steels are often coated with zinc, aluminum, or terne to improve their resistance to corrosion. The coating may be applied by submerging the steel in a container filled with the molten metal (hot-dipped) or by electroplating. The formability of these coated steels is the same as the formability of the base metal to which they are applied, modified by the frictional effects and handling problems imposed by the coating. The base metal for electrogalvanized and terne-coated steels is normally processed in the same manner as uncoated steel base metal, and the mechanical properties are the same. However, zinc and aluminum hot-dipped coated steels are processed in continuous annealing lines, and this treatment affects the mechanical properties and the formability of the base metal. Conventional commercial-quality (CQ), drawing-quality (DQ), and drawing-quality special-killed (DQSK) hot-dipped coated steels will exhibit higher yield and tensile strengths, higher hardness, lower elongation, and lower r values than uncoated low-carbon steels, and they will generally have poorer formability. Decarburized hot-dipped coated steels (sometimes referred to as IF or DDQSK grades) are not similarly affected by the continuous annealing process, and these grades usually have somewhat better mechanical properties and formability than uncoated low-carbon steels. The coatings can be slippery or abrasive. A terne coating is often applied on difficult parts because of its excellent lubricity. Because many hot-dipped zinc coatings are slippery, they may require slightly more blank holder pressure during press forming than do uncoated steels. Aluminum coatings and electrogalvanized coatings tend to develop higher friction during forming and thus generally require better lubrication than uncoated steels. All steel coatings are softer than the steel base metal and can be scraped or gouged off the base metal surface with sharp burrs on blank edges or rough areas in the processing equipment. Coating that is scraped off tends to build up on the tooling (flaking), thereby producing a poor surface on the formed part. The solution is to maintain sharp cutting edges on blanking tools to minimize burr height, as well as a polished surface in die contact areas. Flaking can also be controlled with improved lubrication practice. Gray cast iron dies, in particular, cause flaking problems. Ideally, cast tooling that is to be used to press form a coated steel should be made of cast steel or nodular cast iron. However, flaking problems in gray cast iron tooling can be eliminated by chrome plating or ion nitriding the tool contact surfaces.

Selection of Steel Sheet Steel sheet selection should be based on an understanding of available grades of sheet and forming requirements. Other factors that should be considered when selecting a material for forming into a particular part include: • • • • • • • • • • • •

Purpose of the part and its service requirements Thickness of the sheet metal and allowable tolerances Size and shape of blanks for the forming operation Equipment available for forming Quantities required Available handling equipment for sheets or coils Local availability of sheet products Surface characteristics of the steel sheet Special finishes or coatings for appearance or for corrosion resistance Aging propensity and its relation to time before use Strength of the steel sheets as-delivered Strength requirements in the formed part

Because these factors may be interdependent and large quantities are generally used in part manufacture, it is often desirable that steel selection be made after consultation with either the technical representatives of suppliers or the steel producer. Some parts require specialized low-carbon steel that has been processed to enhance a given mechanical property. These are other less critical formed parts that can use a wide selection of both hot-rolled and cold-rolled steel sheet. The user and producer should understand not only how steels are produced but what a steel mill can do to obtain specific properties in order to prevent the purchase of a steel possessing unwanted properties. The user should be aware of special steels that, although more costly, may reduce production costs and forming problems, resulting in per-part savings. Low-carbon steels, coated and uncoated, are generally supplied as commercial-quality, drawing-quality, and drawing-quality special-killed grades. Some steel mills also offer specialized grades, such as interstitial-free deep-drawing steels and enameling steels. Some of the forming characteristics of the more commonly used formable grades are: • Commercial quality: Available in hot-rolled, cold-rolled, and coated grades. The least expensive grade of sheet steel. Subject to aging (mechanical properties may deteriorate with time). Not intended for difficult-to-form shapes • Drawing quality: Available in hot-rolled, cold-rolled, and coated grades. Exhibits better ductility than CQ grade steels, but has low r values. Subject to aging (mechanical properties may deteriorate with time). Has excellent base metal surface quality


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• Drawing-quality special-killed: Available in hot-rolled, cold-rolled, and coated grades, with good forming capabilities. Not subject to aging (mechanical properties do not change with time) • Interstitial-free steels: Available in cold-rolled and coated grades, with excellent forming capabilities. Not subject to aging (mechanical properties do not change with time) • Enameling steels: Available in cold-rolled grades. Various types of processing are used to make a product that is satisfactory for porcelain enameling. All grades have good forming capabilities • Higher-strength steel sheets: Available in hot-rolled, cold-rolled, and coated grades. Various types of processing are used to obtain the desired strength levels. In general, the formability of these grades decreases as yield strength increases. Springback may be a problem at lower sheet thicknesses Steel sheet selection can be assisted by circle grid analysis, which provides a reliable description of the strain condition in press-formed shapes and which indicates whether the steel is capable of making the required shape or whether a more formable grade is required. Also, if a circle grid analysis shows severe strains with good-quality sheet steel that has normal mechanical properties, this is a strong indication that some modifications will have to be made in the forming process if production is to be maintained. By accurately identifying the problem areas in sheet metal forming operations, circle grid analysis can produce significant savings for both the manufacturers and the users of steel sheet. REFERENCES 1. S.P. Keeler, Understanding Sheet Metal Formability, Machinery, Vol 74 (No. 6−11), Feb-July 1968 2. S.P. Keeler and W.G. Brazier, Relationship Between Laboratory Material Characterization and Press-Shop Formability, in Microalloying 75, Union Carbide Corporation, 1977, p 517−528 3. N.J. Petch, The Ductile-Cleavage Transition in Alpha-Iron, Fracture, B.L. Averback et al., Ed., Technology Press, 1959 4. D.J. Blickwede, Micrometallurgy by the Millions, Trans. ASM, Vol 61, 1968, p 653−679 SELECTED REFERENCES • W.F Hosford and R.M. Caddell, Metal Forming⎯Mechanics and Metallurgy, Prentice-Hall, 1983 • G. Sachs and H.E. Voegeli, Principles and Methods of Sheet Metal Fabricating, Reinhold Publishing, 1966 • "Sheet Steel Formability," Committee of Sheet Steel Producers, American Iron and Steel Institute, 1984


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ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Fabrication Characteristics of Carbon and Low-Alloy Steels Bulk Formability of Steels BULK FORMABILITY, also known as workability, refers to the relative ease with which a metal can be shaped through deformation processes such as forging, extrusion, or rolling. Bulk formability is related to sheet formability in only the broadest sense, in that both characteristics provide quantitative estimates of the strength and ductility of a metal. The latter property⎯ductility, or the resistance of the material to failure⎯is usually of primary concern in describing both bulk and sheet formability.

Formability Characteristics Bulk Versus Sheet Formability. To clarify the distinction between bulk formability and sheet formability, it may be useful to compare and contrast the types of deformation that occur during typical bulk and sheet forming processes. In both processes, the surfaces of the deforming metal are in contact with forming tools, and friction may have a major influence on material flow. In bulk forming, the surface-to-volume ratio of the formed part increases considerably under the action of largely compressive stresses. Plastic deformation is much more prevalent than elastic deformation; therefore, elastic recovery after deformation is negligible. Important material characteristics include flow stress, failure behavior, and the metallurgical transformations that characterize the alloy system in question. In sheet forming operations, the metal is plastically deformed by tensile loads, often without significant changes in sheet thickness or surface characteristics. The magnitudes of plastic and elastic deformation may be similar, resulting in a significant amount of elastic recovery or springback. A key characteristic of a material is the plastic-strain ratio r, the resistance of the sheet to thinning during deformation. The major emphasis in determining both the bulk formability and sheet formability of materials is on measuring and predicting the limits of deformation before fracture. A useful tool for graphically depicting the bulk formability of a material is the workability diagram, which indicates the locus of normal free-surface strains that result in fracture. The analogous concept for sheet formability is the forming limit diagram. Tests for Bulk Formability. Both of these graphic depictions of formability rely on data gathered from laboratory formability tests. A wide variety of tests are used to determine bulk formability, ranging from general tests (tension and torsion tests, for example) to specialized tests that have a very narrow scope and range of application. Test procedures commonly used for determining the bulk formability of steels are covered in the section "Formability Tests" in this article. Bulk Formability of Carbon and Alloy Steels. Despite the large number of available compositions, all the materials in this category exhibit essentially similar bulk formability characteristics. Exceptions to this are steels containing free-machining additives such as sulfides; these materials are not as receptive to bulk forming as nonfree-machining grades (Fig. 1 ). Fig. 1 Comparison of the bulk formability of carbon and low-alloy steels with the formability of resulfurized grades. TM is the absolute melting temperature of the alloys. Source: Ref 1

Generally, the bulk formability of carbon and alloy steels improves as the deformation rate increases. The improvement has been primarily attributed to the increased heat of deformation generated at high deformation rates. Because steels are the most commonly forged materials, a particularly important aspect of their bulk formability is forgeability, the ability to flow readily and fill forging die recesses without fracturing. An important measure of forgeability is flow stress, the amount of force required to deform the material at a specific temperature and strain rate. The section "Evaluating Forgeability" in


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this article includes information on measuring flow stress. This article will present procedures for various formability tests used for carbon and alloy steels. The metallurgy and thermomechanical processing of high-strength low-alloy (microalloyed) steels will also be discussed.

Formability Tests Tests for bulk formability can be divided into two main categories: primary tests and specialized tests. Primary tests such as tension and torsion tests have somewhat limited utility; specialized tests more closely simulate the deformation experienced in actual bulk forming processes and may give better indications of formability. Primary Tests The primary tests for workability are those for which the stress state is well known and controlled. Generally, these are small-scale laboratory simulation tests. The tension test is widely used to determine the mechanical properties of a material (Ref 1). Uniform elongation, total elongation, and reduction in area at fracture are frequently used as indices of ductility. However, the extent of deformation possible in a tension test is limited by the formation of a necked region in the tension specimen. This introduces a triaxial tensile stress state and leads to fracture. For carbon and alloy steels, tension tests are primarily used under special high strain rate, hot tension test conditions to establish the range of hot-working temperatures. The principal advantage of hot tension testing for carbon and alloy steels is that minimum and maximum hot-working temperatures are clearly established. Most commercial hot tensile testing is done with a Gleeble unit, which is a high strain rate, high-temperature testing machine (Ref 2). A solid buttonhead specimen that has a reduced diameter of 6.35 mm (0.250 in.) and an overall length of 88.9 mm (3.5 in.) is held horizontally by water-cooled copper jaws (grips), through which electric power is introduced to resistance heat the test specimen (Fig. 2 ). Specimen temperature is monitored by a thermocouple welded to the specimen surface at the middle of its length. The thermocouple, with a function generator, controls the heat fed into the specimen according to a programmed cycle. Therefore, a specimen can be tested under time and temperature conditions that simulate hot-working sequences. Fig. 2 Gleeble test unit used for hot tension and compression testing. (a) Specimen in grips showing attached thermocouple wires and linear variable differential transformer (LVDT) for measuring strain. (b) Close-up of a test specimen. Courtesy of Duffers Scientific, Inc.

The specimen is loaded by a pneumatic-hydraulic system. The load can be applied at any time in the thermal cycle. Temperature, load, and crosshead displacement are measured as a function of time. The percent reduction in area is the primary result obtained from the hot tension test. This measure of ductility is used to assess the ability of the material to withstand crack propagation. Reduction in area adequately detects small ductility variations in materials caused by composition or processing when the material is of low-to-moderate ductility. It does not reveal small ductility variations in materials of very high ductility. In the torsion test, deformation is caused by pure shear, and large strains can be achieved without the limitations imposed by necking (Ref 3). Because the strain rate is proportional to rotational speed, high strain rates are easily obtained. Moreover, friction has no effect on the test, as it does in compression testing. The stress state in torsion may represent the typical stress in metalworking processes, but deformation in the torsion test is not an accurate simulation of metalworking processes, because of excessive material reorientation at large strains. Fracture data from torsion tests are usually reported in terms of the number of twists to failure or the surface fracture strain to failure. Figure 3 shows the relative hot workability of two AISI carbon and alloy steels as indicated by the torsion test. The test identifies the optimal hot-working temperature for each of the two steels. The section "Evaluating Forgeability" in this article contains information on the use of torsion testing to evaluate the forgeability of carbon and alloy steels.


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Fig. 3 Ductility of two AISI carbon and alloy steels determined in hot torsion tests. Arrows denote suitable hot-working temperatures.

The compression test, in which a cylindrical specimen is upset into a flat pancake, is usually considered to be a standard bulk formability test. The average stress state during testing is similar to that in many bulk deformation processes, without introducing the problems of necking (in tension) or material reorientation (in torsion). Therefore, a large amount of deformation can be achieved before fracture occurs. The stress state can be varied over wide limits by controlling the barreling of the specimen through variations in geometry and by reducing friction between the specimen ends and the anvil with lubricants. Compression testing has developed into a highly sophisticated test for formability in cold upset forging, and it is a common quality control test in the hot forging of carbon and alloy steels. Compression forging is a useful method of assessing the frictional conditions in hot working. The principal disadvantage of the compression test is that tests at a constant, true strain rate require special equipment. Ductility Testing. The basic hot ductility test consists of compressing a series of cylindrical or square specimens to various thicknesses or to the same thickness with varying specimen length-to-diameter (length-to-width) ratios. The limit for compression without failure by radial or peripheral cracking is considered to be a measure of bulk formability. This type of test has been widely used in the forging industry. Longitudinal notches are sometimes machined into the specimens before compression. Because the notches apparently cause more severe stress concentrations, they enable the test to provide a more reliable index of the workability to be expected in a complex forging operation. The bend test is useful for assessing the formability of thick steel sheet and plate. Generally, this test is most applicable to cold-working operations. The principal stress and strains developed during bending are defined in Fig. 4 . The critical parameter is the width-to-thickness ratio, w/t. If w/t > 8, bending occurs under plane-strain conditions (ε2 = 0) and σ2/σ1 = 0.5. If w/t > 8, the bend ductility is independent of the exact w/t ratio. If w/t < 8, then the stress state and bend ductility depends strongly on the width-to-thickness ratio. Fig. 4 Bend region defining the direction of principal stresses and strains in bend testing

Specialized Tests In the plane-strain compression test, the specimen is a thin plate or sheet that is compressed across the width of the strip by narrow platens that are wider than the strip. The elastic constraints of the undeformed shoulders of material on each side


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of the platens prevent extension of the strip in the width dimension; hence the term plane strain. Deformation occurs in the direction of platen motion and in the direction normal to the length of the platen. To ensure that lateral spread is negligible, the width of the strip should be at least six to ten times the breadth of the platens. To ensure that deformation beneath the platens is essentially homogeneous, the ratio of platen breadth to strip thickness, b/t, should be between 2 and 4 at all times. It may be necessary to change the platens during testing to maintain this condition. True strains of 2 can be achieved by carrying out the test in increments in order to provide good lubrication and to maintain the proper b/t ratio. Because of its geometry, this test is more applicable to rolling operations than to forging. The partial-width indentation test is similar to the plane-strain compression test, but it does not subject the test specimen to true plane-strain conditions (Ref 4). In this test, a simple slab-shaped specimen is deformed over part of its width by two opposing rectangular anvils having widths smaller than that of the specimen. Upon penetrating the workpiece, the anvils longitudinally displace metal from the center, creating over-hangs (ribs) that are subjected to secondary, nearly uniaxial tensile straining. The material ductility under these conditions is indicated by the reduction in the rib height at fracture. The test geometry has been standardized. One advantage of this test is that it uses a specimen of simple shape; another is that as-cast materials can be readily tested. One edge of the specimen can contain original surface defects. The test can be conducted hot or cold. The secondary-tension test is a modification of the partial-width indentation test. In this test, a hole or a slot is machined in the slab-type specimen adjacent to where the anvils indent the specimen. With this design, the ribs are sufficiently stretched to ensure fracture in even the most ductile materials. The fracture strain is based on reduction in area where the rib is cut out so that the fracture area can be photographed or traced on an optical comparator. Ring Compression Test. When a flat ring-shaped specimen is upset in the axial direction, the resulting change in shape depends only on the amount of compression in the thickness direction and the frictional conditions at the die/ring interfaces. If the interfacial friction were zero, the ring would deform in the same manner as a solid disk, with each element flowing outward radially at a rate proportional to its distance from the center. In the case of small, but finite, interfacial friction, the outside diameter is smaller than in the zero-friction case. If the friction exceeds a critical value, frictional resistance to outward flow becomes so high that some of the ring material flows inward to the center. Measurements of the inside diameters of compressed rings provide a particularly sensitive means of studying interfacial friction because the inside diameter increases if the friction is low and decreases if the friction is higher. The ring test, then, is a compression test with a built-in frictional measurement. Therefore, it is possible to measure the ring dimensions and compute both the friction value and the basic flow stress of the ring material at the strain under the given deformation conditions. The ring compression test, can be used to measure the flow stress under high-strain practical forming conditions. The only instrumentation required is that for measuring the force needed to produce the reduction in height.

Evaluating Forgeability The hot forging of carbon and alloy steels into intricate shapes is rarely limited by forgeability aspects, with the exception of the free-machining grades mentioned earlier. Section thickness, shape complexity, and forging size are limited primarily by the cooling that occurs when the heated workpiece comes into contact with the cold dies. For this reason, equipment that has relatively short die contact times, such as hammers, is often preferred for forging intricate shapes in steel. Because forging is a complex process, a single test cannot be relied on to determine forgeability. However, several testing techniques have been developed for predicting forgeability, depending on alloy type, microstructure, die geometry, and process variables. This section will summarize some of the common tests for determining formability in open-die and closed-die forgings. Forgeability Hot Twist Testing. One common means of measuring the forgeability of steels is the hot twist test. As the name implies, this test involves twisting heated bar specimens to fracture at a number of different temperatures selected to cover the possible hot-working temperature range of the test material. The number of twists to fracture and the torque required to maintain twisting at a constant rate are reported. The temperature at which the number of twists is the greatest, if such a maximum exists, is assumed to be the optimal hot-working temperature of the test material. Figure 5 shows the forgeabilities of several carbon steels as determined by hot twist testing. More information on the hot twist test is available in Ref 5, 6, and 7. Fig. 5 Forgeabilities of various carbon steels as determined using hot twist testing. Source: Ref 5


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Wedge-Forging Test. In this test, a wedge-shaped piece of metal is machined from a cast ingot or wrought billet and forged between flat, parallel dies (Fig. 6 ). The dimensions of the wedge must be selected so that a representative structure of the ingot is tested. Coarse-grain materials require larger specimens than fine-grain materials. The wedge-forging test is a gradient test in which the degree of deformation varies from a large amount at the thick end (h2) to a small amount or no deformation at the thin end (h1). The specimen should be used on the actual forging equipment in which production will occur to allow for the effects of deformation velocity and die chill on workability. Fig. 6 Specimen for the wedge-forging test. (a) as-machined specimen. (b) Specimen after forging

Tests can be made at a series of preheat temperatures, beginning at about nine-tenths of the solidus temperature or the incipient melting temperature. After testing at each temperature, the deformation that causes cracking can be established. In addition, the extent of recrystallization as a function of strain and temperature can be determined by performing metallographic examinations in the direction of the strain gradient. The sidepressing test consists of compressing a cylindrical bar between flat, parallel dies where the axis of the cylinder is parallel to the surfaces of the dies. Because the cylinder is compressed on its side, this testing procedure is termed sidepressing. This test is sensitive to surface-related cracking and to the general unsoundness of the bar because high tensile stresses are created at the center of the cylinder. For a cylindrical bar deformed against flat dies, the tensile stress is greatest at the start of deformation and decreases as the bar assumes more of a rectangular cross section. The degree of tensile stress can be reduced at the outset of the tests by changing from flat dies to curved dies that support the bar around part of its circumference. The typical sidepressing test is conducted with unconstrained ends. In this case, failure occurs by ductile fracture on the expanding end faces. If the bar is constrained to deform in plane strain by preventing the ends from expanding, deformation will be in pure shear, and cracking will be less likely. Plain-strain conditions can be achieved if the ends are blocked from longitudinal expansion by machining a channel or cavity into the lower die block. The notched-bar upset test is similar to the conventional upset test, except that axial notches are machined into the test specimens. The notched-bar test is used with materials of marginal forgeability for which the standard upset test may indicate an erroneously high degree of workability. The introduction of notches produces high local stresses that induce fracture. The high levels of tensile stress in the test are believed to be more typical of those occurring in actual forging operations.


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Test specimens are prepared by longitudinally quartering a forging billet, thus exposing center material along one corner of each test specimen. Notches with 1.0 or 0.25 mm (0.04 or 0.01 in.) radii are machined into the faces. A weld button is frequently placed on one corner to identify the center and surface material of alloys that are difficult to forge because of segregation. Specimens are heated to predetermined temperatures and upset about 75%. The specimen is oriented with the grooves (notches) in the vertical direction. Because of the stress concentration effect, ruptures are most likely to occur in the notched areas. These ruptures can be classified according to the rating system shown in Fig. 7 . A rating of 0 indicates that no ruptures are observed, and higher numbers indicate an increasing frequency and depth of rupture. Fig. 7 Suggested rating system for notched-bar upset test specimens that exhibit progressively poorer forgeability. A rating of 0 indicates freedom from ruptures in the notched area.

Truncated-Cone Indentation Test. This test involves the indentation of a cylindrical specimen by a conical tool. As a result of the indentation, cracking is made to occur beneath the surface of the testpiece at the tool/material interface. The reduction (measured at the specimen axis) at which cracking occurs can be used to compare the workability of different materials. Alternatively, the reduction (stroke) at which a fixed crack width is produced or the width of the crack at a given reduction can be used as a measure of workability. The truncated cone was developed as a test that minimizes the effects of surface flaws and the variability they produce in workability (Ref 8). This test has been primarily used in cold forging. Flow Localization Complex forgings frequently develop regions of highly localized deformation. Shear bands may span the entire cross section of a forging and, in extreme cases, produce shear cracking. Flow localization can arise from constrained deformation due to die chill or high friction. However, flow localization can also occur in the absence of these effects if the metal undergoes flow softening or negative strain hardening. Nonisothermal Upset Test. The simplest workability test for detecting the influence of heat transfer (die chilling) on flow localization is the nonisothermal upset test, in which the dies are much colder than the workpiece. Zones of flow localization must be made visible by sectioning and metallograhic preparation. The sidepressing test conducted in a nonisothermal manner can also be used to detect flow localization. Several test specimens are sidepressed between flat dies at several workpiece temperatures, die temperatures, and working speeds. The formation of shear bands is determined by metallography. Flow localization by shear band formation is more likely in the sidepressing test than in the upset test. This is due to the absence of a well-defined axisymmetric chill zone. In the sidepressing of round bars, the contact area starts out as 0 and builds up slowly with deformation. In addition, because the deformation is basically plane strain, surfaces of zero extension are present, along which block shearing can initiate and propagate. These are natural surfaces along which shear strain can concentrate into shear bands. Flow Stress and Forging Pressure


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Flow stresses and forging pressures can be obtained from torque curves generated in hot twist tests or from hot compression or tension testing. Figure 8 shows torque versus temperature curves for several carbon and alloy steels obtained from hot twist testing. These data show that the relative forging pressure requirements for this group of alloys do not vary widely at normal hot-forging temperatures. A curve for AISI type 304 stainless steel is included to illustrate the effect of higher alloy content on flow strength. Fig. 8 Deformation resistance versus temperature for various carbon and alloy steels. Source: Ref 9

Figure 9 shows actual forging pressure measurements for 1020 and 4340 steels and AISI A6 tool steel for reductions of 10 and 50%. Forging pressures for 1020 and 4340 vary only slightly at identical temperatures and strain rates. Considerably greater pressures are required for the more highly alloyed A6 material, and this alloy also exhibits a more significant increase in forging pressure with increasing reduction. Fig. 9 Forging pressure versus temperature for three steels. Data are shown for reductions of 10 and 50%; strain rate was constant at 0.7 s−1 Source: Ref 10

Flow Stress in Compression. Ideally, the determination of flow stress in compression should be carried out under


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isothermal conditions (no die chilling) at a constant strain rate and with a minimum of friction in order to minimize barreling. These conditions can be met with conventional servohydraulic testing machines. In flow stress determination, a specific load is applied to a cylindrical specimen of known height, the load is removed, and the new height is determined in order to calculate a true strain. Upon relubrication, the specimen is subjected to an increased load, unloaded, and measured. The cycle is then repeated.

Microalloyed Steels* [*Portions of this section are adapted from S.S. Hansen, Microalloyed Plate and Bar Products: Production Technology, in Fundamentals of Microalloying Forging Steels, The Metallurgical Society, 1987, p 155−172] Microalloying⎯the use of small amounts of elements such as vanadium and niobium to strengthen steels⎯has been in practice since the 1960s to control the microstructure and properties of low-carbon steels. Most of the early developments were related to plate and sheet products in which microalloy precipitation, controlled rolling, and modern steelmaking technology were combined to increase strength significantly relative to that of other low-carbon steels. The application of microalloying technology to forging steels has lagged behind that of flat-rolled products because of the different property requirements and thermomechanical processing of forging steels. Forging steels are commonly used in applications in which high strength, fatigue resistance, and wear resistance are required. These requirements are most often filled by medium-carbon steels. Thus, the development of microalloyed forging steels has centered on grades containing 0.30 to 0.50% C. Regardless of product form (plate, bar, or forging), microalloyed steels are a classic example of a successful metallurgical innovation in which alloying additions and thermomechanical processing have been brought together effectively to attain improved combinations of engineering properties through microstructural control. This practice is relatively inexpensive because only small concentrations of the alloying elements (typically niobium, vanadium, or titanium) are needed to form carbides or carbonitrides. Where possible, the associated thermomechanical processing is introduced merely as a modification of the final hot-working operation. Although precipitation hardening with fine carbonitrides and substructural changes due to warm rolling of austenite-ferrite mixtures can contribute to the strengthening of microalloyed steels, the microstructural feature that ultimately provides a favorable balance of strength and toughness is a small ferritic grain size. The following sections compare the processing of microalloyed plate and bar products. More information on the metallurgy and properties of microalloyed steel forgings is available in the article "High-Strength Low-Alloy Steel Forgings" in this Volume. Processing of Microalloyed Plate Steels Figure 10 shows a temperature-time profile for the rolling of 19 mm (3=4in.) microalloyed steel plate. Initially, slabs are reheated to temperatures in the range of 1100 to 1250 °C (2010 to 2280 °F). The rolling operation itself generally involves two distinct stages: high-temperature rolling or roughing and a lower-temperature series of deformation steps designated as finishing. If the roughing and finishing operations are continuous, the process is termed hot rolling, but if there is a delay between the two stages, as shown in Fig. 10 , the process is referred to as controlled rolling. After rolling, the plate is usually air cooled, although a recently developed technology involves water spray cooling of the plate after finish rolling. Fig. 10 Temperature-time profile for controlled rolling of 19 mm (3=4in.) thick microalloyed steel plate. TR, recrystallization temperature


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Overall, the plate-rolling operation lends itself to considerable control of the thermomechanical treatment. The slab reheat temperature can be reduced if desired. In fact, some rolling strategies involve only reheating to 960 °C (1760 °F) prior to rolling. Delays can be built into the rolling operation (although with some penalty in productivity), and a considerable range of finishing temperatures can be achieved. This operation can accommodate the most severe controlled-rolling schedules, including the deformation of austenite-ferrite mixtures. In metallurgical terms, the controlled-rolling operation in microalloyed steels serves two purposes. The first is to refine the relatively coarse, as-reheated austenitic microstructure by a series of high-temperature rolling and recrystallization steps. The second purpose of the rolling operation is to impose a moderate-to-heavy reduction in a temperature regime where austenite recrystallization is inhibited between rolling passes (below the recrystallization temperature indicated in Fig. 10 ) such that the plastically deformed austenite grains remain pancaked. Subsequent transformation after rolling into ferritic microstructures results in the desired fine grain size and associated mechanical properties. The various parts of the rolling operation are discussed separately and in greater detail in the following sections of this article. Thermomechanical Treatment (Rolling). Typically, the initial rolling passes are conducted at relatively high temperatures, just below the slab-reheating temperature. At these temperatures, each deformation step is usually followed by rapid recrystallization and grain growth. Recently, a thermomechanical processing procedure called recrystallization controlled rolling has been proposed. It combines repeated deformation and recrystallization steps with the addition of austenite grain-growth inhibitors such as titanium nitride to refine the starting austenitic grain size and to restrict grain growth after recrystallization. Such processing would obviate the need for low-temperature controlled rolling (Ref 11, 12, 13). However, even with optimum compositions and the adoption of rather difficult reduction schedules, there seems to be a limit to the degree of austenitic refinement that can be achieved by repeated recrystallization; the finest recrystallized austenitic grain sizes produced by this process are about 15 µm (600 µin.). Depending on the subsequent cooling rate, transformation can then result in a ferritic grain size of 6 to 8 µm (240 to 320 µin.) (Ref 13). This is useful degree of structural refinement and is appropriate in those cases where controlled rolling is not possible (for example, due to mill load constraints). However, still finer grain sizes can be attained through the use of additional microalloying elements along with a controlled-rolling sequence in which austenite recrystallization is substantially retarded during the later rolling passes. This process develops a pancaked grain morphology and a much higher surface area per unit volume than are possible in recrystallized austenite (Ref 14). During this process, the austenite recrystallization and carbonitride precipitation reactions are coupled in the sense that each is greatly influenced by the other. Transformation to Ferrite. The transformation of the austenite grain into ferritic microstructures determine the final grain size and associated mechanical properties of the microalloyed plate. The effects of austenitic morphology and the transformation temperature range (as governed by alloy content, rolling deformation, and cooling rate) are of the greatest importance. Even after the minimum austenitic grain thickness has been produced, the temperature range of the austenite-to-ferrite transformation must be controlled to determine the reaction kinetics. Increasing the ferritic nucleation rate and decreasing the ferritic growth rate can produce a finer ferritic grain size. These effects are generally achieved by alloying or controlled cooling. Precipitation and Substructural Strengthening. Although grain refinement offers the best combination of strength and toughness, there is a practical limit to the yield strength level that can be achieved with this strengthening mechanism alone: about 450 MPa (65 ksi) for a grain size of 3 µm (120 µin.). For higher strength levels, additional strengthening mechanisms must be used; however, these mechanisms can have deleterious effects on toughness (Ref 15). For example, as vanadium is added to a controlled-rolled niobium-containing microalloyed plate steel, the yield strength is increased about 7 MPa (1 ksi) for every 0.01% increase in vanadium. This strengthening is due to the precipitation of vanadium-rich carbonitrides in the ferrite during air cooling of the plate after rolling. For a typical vanadium content of 0.10%, this precipitation strengthening can raise the yield strength to about 525 MPa (76 ksi). This results in the deformation of austenite-ferrite mixtures and the development of a warm-worked ferritic substructure. As shown in Fig. 11 , strength increases progressively as rolling is continued to lower temperatures. At the same time, toughness is also reduced; therefore, this strengthening mechanism is used in microalloyed steels mainly as a way to reach the last increment of required strength. Fig. 11 Yield strength of a microalloyed steel as a function of finishing temperature. Grain size: 5 µm (200 µin.). Source: Ref 16


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High-Strength Low-Alloy (HSLA) Plate Products. The basic metallurgical principles discussed above are used to produce a range of as-rolled microalloyed plate steels at thicknesses to 102 mm (4 in.) and yield strengths as high as 550 MPa (80 ksi). In these steels, niobium is added to control the recrystallization temperature, manganese is often used to control austenite formation, and vanadium is introduced for precipitation strengthening. Finish rolling is carried out at temperatures low enough to produce substructural strengthening. Processing of Microalloyed Bars In contrast to the plate-rolling process, the thermomechanical treatment possible on a modern bar mill (Fig. 12 ) is somewhat limited in scope. For example, the temperature-time profile for the rolling of 44 mm (13=4in.) diam bar shown in Fig. 13 can be compared to the plate-rolling profile shown in Fig. 10 . There are clear differences between these rolling processes: • Lower reheat temperatures, typically in the range of 1100 to 1200 °C (2010 to 2190 °F), are used in bar rolling. This lower temperature, in combination with the generally higher carbon levels in bar products, limits the amount of niobium that can be dissolved upon reheating. For example, in a 0.20% C steel, only about 0.01% Nb is soluble at 1100 °C (2010 °F). In contrast, vanadium is still readily soluble at bar reheat temperatures. Consequently, in HSLA bar grades, vanadium is the microalloying element commonly used to obtain to the highest possible strength levels • Even though the lower reheat temperatures typical of bar products place some limitations on the use of different microalloying elements, these lower temperatures do provide for a finer as-reheated austenitic grain size than is typical of slabs reheated for conversion to plate. With a small titanium addition and continuous casting, as-reheated austenitic grain sizes of 50 to 60 µm (0.0020 to 0.0024 in.) can be achieved in billets destined for bar • Finishing temperatures in bar rolling are relatively high, even with the use of interstand cooling Consequently, recrystallization controlled rolling becomes quite important in bar rolling, and the rolling strategy must be designed to produce the finest possible recrystallized austenitic grain size. Subsequent control of the austenite-to-ferrite transformation range is still important to maximize ferritic grain refinement. Nevertheless, as discussed earlier for plate, the ferritic grain size that can be produced on transformation from a recrystallized austenite is limited compared to the grain size that can be produced on transformation from austenitic grains that have been flattened by rolling below the recrystallization temperature. Thus, while moderate grain refinement can be achieved in an as-rolled microalloyed bar, this grain size will be somewhat coarser than the grain size of controlled-rolled HSLA plates. Fig. 12 Controlled-rolling process for microalloyed steel bar. Source: Ref 17

Fig. 13 Temperature-time profile for the controlled rolling of 44 mm (13=4in.) diam microalloyed steel bar. Compare with Fig. 10 .


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Alternative Strengthening Mechanisms in Microalloyed Bar Steels. Because the degree of ferritic grain refinement possible in as-rolled microalloyed bar steels is somewhat limited, and because substructural strengthening is not possible, alternative strengthening mechanisms must be employed to reach yield strength levels comparable to those of plate grades. For example, in the alloy design of microalloyed bar steels, precipitation and pearlite strengthening must be relied on to a greater extent than in the design of plates. In view of the limited solubility of niobium or titanium at the reheat temperatures used in bar processing, vanadium is usually used to obtain the required level of precipitation strengthening in HSLA bar grades. Precipitation of V(C,N) during or after transformation can provide significant strengthening increments. In this regard, nitrogen level is also of importance. Judicious selection of both the vanadium and nitrogen levels is required to produce the desired level of precipitation strengthening. Similarly, an increase in the carbon content and thus the pearlite volume fraction of a bar steel can also be used to increase strength (Ref 18, 19, 20). In addition to moderate grain refinement, precipitation hardening with VN and an increase in the pearlite volume fraction can be used to produce yield strength levels up to 625 MPa (91 ksi) in microalloyed bars. Of course, these two strengthening mechanisms have very deleterious effects on toughness. High-Strength Low-Alloy Bar Products. The alloy and process design principles discussed above are employed to produce a reasonable selection of as-rolled microalloyed bar steels. Yield strength levels up to 625 MPa (91 ksi) have been publicized, although the available yield strength level is influenced by bar thickness. In these steels, titanium is sometimes added to control austenitic grain growth (on reheating and after recrystallization during rolling), carbon and manganese can be balanced to control the transformation temperature range, vanadium and nitrogen are used for precipitation strengthening, and carbon is increased (as required) to raise the pearlite fraction of the ferrite and pearlite microstructural aggregate. Compared to the higher-carbon quenched and tempered grades that are currently used in competitive applications, these microalloyed bar steels offer comparable strength at lower cost because less alloy is required and heat-treating costs are eliminated. However, the toughness of the microalloyed bar grades developed to date is still somewhat lower than that of the higher-carbon, heat-treated steels. While the toughness levels currently available in commercial microalloyed bar steels may be adequate for many applications, considerable effort is being made at the present time to improve the toughness of microalloyed bar grades. Comparison of Microalloyed Plate and Bar Products The differing processing approaches for plate and bar are reflected in the microstructure and properties that are ultimately developed in these two product forms. Consider, as an example, the attributes of microalloyed plate and bar grades at the 550 MPa (80 ksi) yield strength level. The bar product (Fig. 14 ) has a coarser ferritic grain size and a significantly higher pearlite volume fraction (due to the higher carbon content) than the plate product. The comparison of compositions and yield-strengthening increments shown in Fig. 15 reflects these differences in microstructure. The most significant yield strength increment in the plate product is due to ferritic grain refinement, while the bar product must rely more on precipitation strengthening. In both products, however, small strengthening contributions are required beyond the grain size and precipitation hardening level to reach a yield strength of 550 MPa (80 ksi). In plate, this increment is obtained by rolling to develop a ferritic substructure, while in bar a small contribution from pearlite is necessary. Of course, these different strengthening mechanisms have an impact on the toughness that can be achieved in these two product forms (Fig. 16 ). Because it has a finer ferritic grain size and relies less on precipitation strengthening, the plate product exhibits significantly better toughness than the bar product. Fig. 14 Grain structure of a microalloyed steel bar product of composition Fe-0.38C-1.18Mn-0.16V-0.018N. Source: Ref 21


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01 Sep 2005

Fig. 15 Composition and yield strength increments of microalloyed plate and bar steels with yield strengths of 550 MPa (80 ksi)

Composition, % Product Plate

C

Mn

Si

Nb

V

N

0.14

1.45

0.25

0.035

0.08

0.012


Page 898


Bar


0.26

1.38

0.20

01 Sep 2005

0.005

0.16

0.018

Fig. 16 Comparison of the Charpy V-notch impact toughness of microalloyed plate and bar steels with yield strengths of 550 MPa (80 ksi)

Processing of Microalloyed Forging Steels The driving force behind the development of microalloyed forging steels has been the need to reduce manufacturing costs. This is accomplished in these materials by means of a simplfied thermomechanical treatment (that is, a controlled cooling following hot forging) that achieves the desired properties without the separate quenching and tempering treatments required by conventional carbon and alloy steels. Control of Properties. In order to realize the full strengthening potential of microalloying additions, it is necessary to use a soaking temperature prior to forging that is high enough to dissolve all vanadium-bearing precipitates. A soaking temperature above 1100 °C (2010 °F) is preferred. Rapid induction heating methods for bar and billet to conventional commercial forging temperatures of 1250 °C (2280 °F) are acceptable and allow sufficient time for the dissolution of the microalloying constituents. Tensile strength decreases slightly as the finish forging temperature is reduced, but there is no significant effect on yield strength. Ductility and toughness show a significant increase with a reduction in finishing temperature; this is due to grain refinement of the austenite and increased ferrite content. Forgers are beginning to utilize this approach to enhancing the toughness of as-forged microalloyed steel; however, low finish forging temperatures are often avoided to minimize die wear. The specified properties of microalloyed forging steels can be achieved over a wide range of finishing temperatures. One of the most important processing factors affecting the properties of as-forged microalloyed steels is the postforging cooling procedure. Increasing the cooling rate generally increases the yield and tensile strength because it enhances grain refinement and precipitation hardening. At high cooling rates, an optimum can be reached; above this rate the strength reduces due to the suppression of precipitation and the introduction of low-temperature transformation products. The optimum cooling rate and maximum hardness are significantly influenced by the alloy and residual element content of the steel. Nevertheless, through control of the steel composition it is possible to ensure that the specified mechanical properties are achieved over a wide range of section sizes and cooling conditions. Properties of Forged Parts. Because of the improved strength and hardness of microalloyed bar steels, it is possible to use them for the production of many forged parts and eliminate the need for subsequent heat treatment. This is particularly true when air cooling is applied to the as-forged parts. Numerous forgings weighing as little as 1 kg (3 lb) to well over 11 kg (25 lb) have been produced by this approach. The parts produced from microalloyed steel include connecting rods and caps, stub yokes, weld yokes, wheel hubs, stabilizer bars, blower shafts, sucker rods, anchor bolts, and U-bolts. An example of the improved properties that can be obtained is shown in Fig. 17 , in which the cross-sectional hardness of an air-cooled microalloyed 1541 forged part is compared with a similar quenched and tempered 1043 part. The hardness is much more uniform in the microalloyed part, and as a result the fatigue life at this location in the part is an order of magnitude greater than that in the quenched and tempered part. Fig. 17 Comparison of the cross-sectional Rockwell C hardness and fatigue strength of microalloyed steel forgings (air-cooled 1541 grade containing 0.18% V) and medium-carbon quenched and tempered steel forgings (1043 grade). Source: 22


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01 Sep 2005

Fatigue test results Stress MPa

ksi

Cycles to failure

520

75

219,700

550

80

113,200

520

75

29,100

550

80

18,400

Air-cooled 1541 with 0.18% V

Quenched and tempered 1043

Effects of Hot Mill Finishing Temperature (Ref 23). In tests of a grade of 1541 steel microalloyed with 0.10% V, higher hardness and tensile strengths were obtained with higher hot mill finishing temperatures (Fig. 18 ). Reduction in area also increased slightly at higher finishing temperatures. Yield strength and percent elongation values did not vary with finishing temperature over the range investigated. The steel finished at the lowest temperature (970 °C, or 1780 °F) had the highest impact strength in subsequent Charpy V-notch testing (Fig. 19 ). Fig. 18 Effect of hot mill finishing temperature on hardness and tensile properties of 1541 steel containing 0.10% V. Source: Ref 23


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Fig. 19 Effect of hot mill finishing temperature on impact properties of 1541 steel containing 0.10% V. Source: Ref 23

Effects of Forging Temperature. To achieve an optimum balance of strength and toughness properties, forged parts produced from microalloyed steel must be air-cooled through the transformation temperature. Slow cooling rates resulting from batch cooling must be avoided. Because the forging process is the final thermal processing step in the production of parts from microalloyed chemistries, it is important that the forging operation be controlled in the same manner that the steelmaker controls the bar-rolling operation. The key forging variables that require process control are the reheating temperature and the postforging cooling rate (Ref 23). REFERENCES 1. A.M. Sabroff, F.W. Boulger, and H.J. Henning, Forging Materials and Practices, Reinhold, 1968 2. E.F. Nippes, W.F. Savage, B.J. Bastian, and R.M. Curran, An Investigation of the Hot Ductility of High-Temperature Alloys, Weld. J., Vol 34, April 1955, p 183−196s 3. M.J. Luton, Hot Torsion Testing, in Workability Testing Techniques, G.E. Dieter, Ed., American Society for Metals, 1984, p 95−133 4. S.M. Woodall and J.A. Schey, Development of New Workability Test Techniques, J. Mech. Work. Technol., Vol 2, 1979, p 367−384 5. Evaluating the Forgeability of Steel, 4th ed., The Timken Company, 1974 6. H.K. Ihrig, The Effect of Various Elements on the Hot Workability of Steel, Trans. AIME, Vol 167, 1946, p 749−777 7. C.L. Clark and J.J. Russ, A Laboratory Evaluation of the Hot Working Characteristics of Metals, Trans. AIME, Vol 167, 1946, p 736−748 8. T. Okamoto, T. Fukuda, and H. Hagita, Material Fracture in Cold Forging⎯Systematic Classification of Working Methods and Types of Cracking in Cold Forging, Sumitomo Search, No. 9, May 1973, p 46; Source Book on Cold Forming, American Society for Metals, 1975, p 216−226 9. C.T. Anderson, R.W. Kimball, and F.R. Cattoir, Effect of Various Elements on the Hot Working Characteristics and Physical Properties of Fe-C Alloys, J. Met., Vol 5 (No. 4), April 1953, p 525−529 10. H.J. Henning, A.M. Sabroff, and F.W. Boulger, "A Study of Forging Variables," Technical Documentary Report ML-TDR-64-95, Battelle Memorial Institute, March 1964 11. W. Roberts, A. Sandberg, T. Siwecki, and T. Werlefors, Prediction of Microstructure Development During Recrystallization Hot Rolling of Ti-V Steels, in HSLA Steels: Technology and Applications, American Society for Metals, 1984, p 67 12. H. Sekine, T. Maruyama, H. Kageyama, and Y. Kawashima, in Thermomechanical Processing of Microalloyed Austenite, The Metallurgical Society, 1982, p 141 13. T. Siwecki, A. Sandberg, W. Roberts, and R. Lagneborg, in Thermomechanical Processing of Microalloyed Austenite, The Metallurgical Society, 1982, p 167


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14. L.J. Cuddy, Metall. Trans. A, Vol 15A, 1984, p 87−98 15. F.B. Pickering, in Microalloying '75, Union Carbide Corporation, 1975, p 9 16. J.H. Little, J.A. Chapman, W.B. Morrison, and B. Mintz, The Microstructure and Design of Alloys, Vol 1, The Metals Society, 1974, p 80 17. T. Sampei, T. Abe, H. Osuzu, and I. Kosazu, Microalloyed Bar for Machine Structural Use, in HSLA Steels: Technology and Applications, American Society for Metals, 1984, p 1063−1070 18. H.J. Kouwenhoven, Trans. ASM, Vol 62, 1969, p 437−446 19. T. Gladman, I.D. McIvor, and F.B. Pickering, J. Iron Steel Inst., Vol 210, 1972, p 916−930 20. F.B. Pickering, in Hardenability Concepts with Applications to Steels, The Metallurgical Society, 1978, p 179 21. B.L. Bramfitt, S.S. Hansen, D.P. Wirick, and W.B. Collins, Development of a Microalloyed Joint Bar, in Microalloyed HSLA Steels, ASM INTERNATIONAL, 1988, p 451−457 22. J.F. Held and B.A. Lauer, Development of Microalloyed Medium Carbon Hot Rolled Bar Products, in HSLA Steels: Technology and Applications, American Society for Metals, 1984, p 1071−1080 23. J.F. Held, Some Factors Influencing the Mechanical Properties of Microalloyed Steel, in Fundamentals of Microalloying Forging Steels, The Metallurgical Society, 1987, p 175−188


Page 902


Machinability of Steels

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ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Fabrication Characteristics of Carbon and Low-Alloy Steels Machinability of Steels Francis W. Boulger, Battelle-Columbus Laboratories (retired) THE MACHINABILITY OF CARBON AND ALLOY STEELS is affected by many factors, such as the composition, microstructure, and strength level of the steel; the feeds, speeds, and depth of cut; and the choice of cutting fluid and cutting tool material. These machining characteristics, in turn, affect the cost of producing steel parts, particularly when the cost of machining represents a major part of the cost of the finished part. This article describes the influence of the various attributes of carbon and alloy steels on machining characteristics. It should be recognized that the relative cost of cutting two steels in a particular operation, such as boring, is not necessarily the same as the relative ease of cutting the same two steels in another operation, such as broaching. Machining processes differ in operational metal removal characteristics; some place a greater premium on high machinability of the workpiece than others. Several common machining processes are listed in approximate decreasing order of machinability requirement, as follows: • • • • • • • • • • • • • •

Internal broaching External broaching Tapping Generation of gear teeth Deep drilling Boring Screw machining with form tools High-speed light-feed screw machining Milling Shallow drilling Planing and shaping Turning with single-point tools Sawing Grinding

The designer's choice of part shape and dimensions largely determines the selection of the machining process. The mechanical properties needed for satisfactory service performance usually dictate the selection of the workpiece material and condition of heat treatment. Consequently, decisions about materials by the manufacturing engineer are generally reduced to choosing between similar grades of steel (for example, between 4140 and 8640) rather than between very different grades (for example, 4140 and 12L14).

Measures of Machinability The term machinability is used to indicate the ease or difficulty with which a material can be machined to the size, shape, and desired surface finish. The terms machinability index and machinability rating are used as qualitative and relative measures of the machinability of a steel under specified conditions. There are no clear-cut or unambiguous meanings for these terms and no standard or universally accepted method of measuring machinability. Historically, machinability judgments have been based on one or more of the following criteria: • Tool life: Measured by the amount of material that can be removed by a standard cutting tool under standard cutting conditions before tool performance becomes unacceptable or tool wear reaches a specified amount • Cutting speed: Measured by the maximum speed at which a standard tool under standard conditions can continue to provide


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satisfactory performance for a specified period • Power consumption: Measured by the power required to remove a unit volume of material under specified machining conditions • Comparisons with a standard steel based on experience in machine shops • Quality of surface finish • Feeds resulting from a constant thrust force Some of the test criteria are best suited to laboratory studies intended to elicit information about the effects of small changes in microstructure, composition, or processing history on machinability. Other types of tests are useful for studying the effects of geometry changes or cutting tool composition. Years ago, machinability ratings were also relied upon as aids for choosing machining conditions to be used on materials unfamiliar to production personnel. This is rarely necessary now because detailed and reliable guides to suitable practices, such as the Machining Data Handbook (Ref 1), are readily available and widely used. Tool life and cutting speed can be related by the equation: VcTn = Ct (Eq 1) where Vc is the cutting speed, T is the tool life, and n and Ct are empirical constants that reflect the cutting conditions under which the tests were made and the machinability of the material. In 1907, Taylor presented Eq 1 to describe single-point turning; the constant Ct is often called the Taylor constant. Because typical values of n for high-speed steel (HSS) tools range from 0.1 to 0.2, small variations in cutting speed are equivalent to enormous changes in tool life. Therefore, it is more practical to measure machinability as the cutting speed necessary to cause tool failure within a specified period, usually 60 min, than as tool life at a specified cutting speed. To determine the machinability of a particular steel, tool life for each of several cutting speeds (with standardized cutting conditions and tool shape) must be measured. Values of n and Ct can be determined from these data, and the cutting speed that corresponds exactly to the specified tool life can be calculated. Tool life tests are used in laboratories to evaluate the effects of changes in tool materials, cutting variables, processing history, or workpiece compositions or microstructure on the ease of removing material. They are also useful for predicting tool life and choosing cutting speeds for industrial operations. There are several criteria that can be used to define the failure of cutting tools. One criterion is the complete destruction of the cutting surface of the tool. A second criterion is the wear of the tool to the extent that the quality of the machined surface becomes unacceptable. Perhaps the most widely used criterion for tool failure is wear of the surface of the tool to some predetermined amount of flank wear. Sometimes, especially in screw machine tests, a specific increase in a part dimension is used to define tool life. Regardless of the criterion adopted for tool failure, any machinability rating that depends on tool life measurements will be affected by the cutting tool. The choice of tool material, the configuration of the tool, the sharpness of the cutting edge, and the efficiency with which the tool is cooled can affect the machinability rating of the steel under test. For example, typical values of n for HSS tools range from 0.1 to 0.2; for carbide tools, typical values of n range from 0.2 to 0.4. Therefore, machinability testing should be carefully standardized (as described in Ref 2 and 3) so that the test reflects the machinability of the material rather than variations in the test procedures. Power Consumption. The forces acting on a tool during cutting, as measured on a dynamometer, can be used to estimate the power consumed in metal cutting. The power consumption (expressed in watts) in cutting operations is approximately equal to the product of the cutting speed, Vc (expressed in units of meters per second), and the component of the cutting force parallel to the cutting direction, Fc (expressed in newtons). In English units, this relation for power consumption is:

(Eq 2) where P is power consumption (at the spindle) in horsepower units, Fc is the cutting force in pounds, and Vc is the cutting speed in feet per minute. To calculate the unit power consumption, which reflects the power requirements for cutting a given quantity of a particular material, it is necessary to divide the power consumption, P, by the metal removal rate, Q (which is typically expressed in units of either cubic centimeters or inches per minute). Therefore, the unit power consumption, Pu, is:

(Eq 3) In cutting operations, Q Á d · f · Vc, where d is the depth of cut and f is the feed rate. In metric units, the unit power consumption, Pu, is given by:


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(Eq 4) where Pu is in joules per cubic centimeter, Fc is in newtons, and f and d are in millimeters. If Fc is given in pounds of force and f and d in inches, then Pu in horsepower per cubic inch per minute is given by:

(Eq 4a) Typical data for unit power consumption of steels machined in different processes are given in Table 1 and Fig. 1 . Note that the unit power consumption increases with increasing hardness, which reflects the resistance of the material to the deformation required in machining operations. Table 1 Average unit power requirements for turning, drilling, and milling plain carbon and alloy steels Unit power, kW/cm3/min (hp/in.3/min)(a)

Hardness, HRC

Turning with a feed of 0.12−0.50 mm/rev (0.005−0.020 in./rev)

Drilling with a feed of 0.05−0.20 mm/rev (0.002−0.008 in./rev)

Milling with a feed of 0.12−0.30 mm/tooth (0.005−0.012 in./tooth)

85−200 HB

0.050−0.064 (1.1−1.4)

0.046−0.059 (1.0−1.3)

0.050−0.064 (1.1−1.4)

35−40

0.064−0.077 (1.4−1.7)

0.064−0.077 (1.4−1.7)

0.068−0.086 (1.5−1.9)

40−50

0.068−0.086 (1.5−1.9)

0.077−0.096 (1.7−2.1)

0.082−0.100 (1.8−2.2)

50−55

0.091−0.114 (2.0−2.5)

0.096−0.118 (2.1−2.6)

0.096−0.118 (2.1−2.6)

55−58 0.155−0.191 (3.4−4.2) 0.118−0.146 (2.6−3.2) (a) Power requirements at spindle-drive motor, corrected for 80% spindle-drive efficiency. Source: Ref 4

0.118−0.146 (2.6−3.2)

Fig. 1 Unit power consumption for surface broaching (HSS tools)

The ranges of values for unit power requirements for a particular hardness level cover the spread for sharp and dull tools. The energy is used in deforming metal in the chips and the surface layers of the workpiece and in overcoming friction. Figure 1 shows that decreasing the amount of metal removed by each tooth on a broach increases the unit power consumption, because the friction between the tool and chip or workpiece increases. The same effects occur when feeds are reduced in turning or milling operations. The choice of cutting tool shape and material and the application of coolants have comparatively little effect on the


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unit power consumption, except by altering the amount of power expended in friction. Unit power consumption cannot be easily correlated with tool life; the factors that affect the unit power consumption are primarily the inherent machinability of the material and the power consumed by friction, while the additional factors that affect tool life include the shape and material of the tool, the temperature of the interface between the tool and chip, and the extent of the abrasive action of the chip on the surface of the tool. Quality of surface finish is another means of assessing the machinability of materials. Because the surface finish of a machined part may affect its performance in service, it is sometimes useful to rate the machining characteristics of candidate materials in terms of the surface finish that can be expected from machining under specified conditions. Such ratings are generally qualitative, although materials that have high machinability ratings, as determined by other rating methods, usually produce smooth surfaces, partly because they are machined at high speeds. Machinability Testing for Screw Machines. Although the principles described have been used for judging the machining characteristics of steels by short-time tests made on small amounts of material and for quantifying various characteristics, they are not useful for setting up screw machines. Therefore, engineers concerned with parts produced on automatic bar machines or with making steels for such applications have developed special procedures for determining machinability (Ref 5). Test parts are produced by simultaneous cutting with form tools and by finish-forming and cutoff operations. The conditions to be used and the records to be kept are described in ASTM E 618. Ratings for different lots are assigned on the basis of maximum production rates for parts meeting specified dimensional and surface roughness tolerances.

Scatter in Machinability Ratings Considerable scatter in the machinability data for the steel chosen as the reference for machinability ratings, B1112, is illustrated in Fig. 2 . Within the composition range permitted for that grade, it was found that unintentional variations in carbon, sulfur, and, principally, silicon contents cause the machinability index of B1112 to vary as much as 20% below or 60% above the nominal value of 100 assigned to it. The data on multiple heats of B1113 indicate a similarly large spread of values for that steel. Fig. 2 Distribution of machinability ratings for B1112 and B1113 steels. Source: Ref 6

The effects of small variations in composition and grain size on machinability are sometimes greater than those from variations in hardness. Therefore, variations in performance should be expected when machining different lots of ostensibly similar material. The scatter in machinability data indicates that the precision of machinability ratings is limited, to some extent at least, by the concept of using the average behavior of a particular grade of steel as the standard for comparison. In general, differences of 5% in machinability ratings are not likely to be significant or reproducible. It has been shown that such a scatter can result from variations in composition that meet the chemical ranges permitted by specifications for the grade (Ref 7).

Machinability Ratings of Steels


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01 Sep 2005

It is generally agreed that machinability, as defined by tool life, depends on or correlates with the following characteristics of the workpiece: • • • • •

Structure Chemical composition Dispersion of second-phase particles Mechanical properties, such as strength and hardness Physical properties, such as thermal conductivity

Variations in these characteristics control the performance of workpieces machined under comparable conditions. Materials with superior machinability give better tool lives at equal cutting speeds or permit higher cutting speeds while maintaining equal tool lives. Adopting either alternative⎯that is, better tool lives or higher cutting speeds⎯improves productivity and lowers machining costs. Most published machinability ratings have been based on the performance of steels in one type of operation (usually turning) and with one type of cutting tool. This is particularly true of data from laboratory tests. Consequently, it is of interest to determine whether data obtained with carbide or coated-carbide tools rank materials in the same order of machinability as ratings obtained with HSS tools. Similarly, it is important to determine whether steels exhibiting superiority in one type of operation, such as turning, will also perform better in drilling, boring, milling, or other operations. In the following two sections, machinability rankings are compared for different cutting tool materials and different machining operations. In both sections, machinability is assessed in terms of cutting speeds. Assuming that the other machining parameters are comparable, the machinability of a material is reflected by its permissible cutting speed for a usefully long tool life. This provides an approach for determining the information mentioned above. Information on suitable cutting speeds, based on experience in many industrial shops, has been collected, evaluated, and published (Ref 1). Usually, the recommended cutting speed gave tool lives of about 2 h for HSS tools and about 1 h for carbide tools. Order of Machinability Rankings With Different Cutting Tool Materials. The cutting speeds used for the cross plots in Fig. 3 are those recommended for turning metals, at cut depths of 1.0 mm (0.04 in.) and at appropriate feeds, with three types of cutting tools. The workpiece materials represented in Fig. 3 include 3 types of stainless steel, 14 grades of constructional steel, and 5 varieties of cast iron. Their hardnesses ranged from 100 to 325 HB. The plots show close and consistent relationships among the turning speeds recommended for the three types of cutting tools. Fig. 3 Correlation among suggested cutting speeds for turning different ferrous metals with indexable-carbide, coated-carbide, and HSS tools. Cutting speeds are from Ref 1.


Page 907



01 Sep 2005

Indexable carbide

HSS

Coated carbide

m/min

sfm

m/min

sfm

m/min

sfm

60

200

240

790

365

1200

64

210

250

820

373

1225

90

300

275

900

411

1350

58

190

221

725

335

1100

106

350

320

1050

457

1500

41

135

150

500

198

650

35

115

142

465

180

600

27

90

134

440

175

575

44

145

174

570

260

850

40

130

160

525

239

785

41

135

150

500

198

650

40

130

161

530

213

700

37

120

155

510

229

750

34

110

130

425

168

550

56

185

221

725

290

950


Page 908



01 Sep 2005

37

120

236

775

290

950

60

200

165

540

213

700

43

140

165

540

213

700

58

190

245

800

305

1000

46

150

198

650

260

850

29

95

114

375

150

500

47

155

189

620

245

800

Table 2 lists the statistical attributes conventionally employed for judging the reliability of the correlations shown in Fig. 3 . The goodness of fit between the points and the trend lines is indicated by the correlation coefficient or r value. For the two plots in Fig. 3 , the r values are high enough to permit the conclusion that the relationships shown in the chart are statistically significant at the 99.9% level. The coefficients of determination (r2 values) indicate the proportion of the total variation in the dependent variable explained by its relationship with the independent variable. High proportions of the variance among permissible cutting speeds (89 and 76%) were accounted for by the relationships between indexable carbides and the other two types of cutting tools. The coefficient of determination for the relationship between suggested cutting speeds for coated carbide and steel tools was smaller and less indicative of a strong correlation. The slope and intercept values in Table 2 describe the mean-square trend lines for the relationships when used in an equation with the form: y = (slope)x + (intercept) Table 2 Correlation data of recommended cutting speeds for turning operations with different types of cutting tool materials The cutting speeds used for these correlation analyses are given in Fig. 3 and are those recommended in Ref 1 for turning 5 types of cast irons, 3 types of stainless steels, and 14 constructional steels. Paired tool materials Correlation statistic Dependent variable

Independent variable

Correlation coefficient

Coefficient of determination

Slope of trend line (Fig. 3 )

Intercept on y-axis (Fig. 3 )

HSS

Indexable carbide

0.8722

0.761

0.35

−34.9

Coated carbide

Indexable carbide

0.9427

0.890

1.41

−3.97

HSS

Coated carbide

0.8387

0.703

3.74

241

The correlation data confirm the opinion that workpiece characteristics that improve the ease of machining a metal will be effective in turning tests made with different types of cutting tools. The trend lines also indicate that changing to a better tool material does not raise the permissible cutting speed an equal or fixed amount for different metals. Order of Machinability Rankings With Different Machining Operations. The cutting speeds prescribed for the same 22 types of ferrous metals were also used to obtain their machinability ratings in different types of operations. The ratings were based on the speeds (in surface feet per minute) recommended for machining with HSS tools (Ref 1). The cutting speeds suggested for 1212 steel, at a hardness level of 150 to 200 HB, were used as the basis for comparison. The machinability rating of a material was taken to be 100 times the ratio of its recommended cutting speed to that for 1212 steel for otherwise comparable cutting conditions. Figure 4 shows a crossplot of the machinability ratings calculated as above for boring and for turning with a cut depth of 1.0 mm (0.04 in.). The correlation chart shows that the agreement was very close between the two types of metal removal operations. Although not shown in the chart, the least-squares trend lines for machinability ratings based on rough reaming and those based on turning at a heavier cut depth (0.38 mm, or 0.15 in.) closely matched the one shown for boring; that is, they fell within the scatter band of points shown for boring. Fig. 4 Correlations among machinability ratings for different materials based on recommended speeds for turning and for boring with HSS tools. See text for details.


Page 909


Turning


01 Sep 2005

Boring

Turning

Boring

95

94

62

56

100

100

57

56

148

141

53

44

90

88

88

85

162

159

57

65

167

165

95

94

68

65

67

65

58

53

90

88

45

41

71

71

35

32

45

44

68

68

74

74

Table 3 gives the correlation statistics comparing the machinability ratings calculated from speeds recommended for different types of operations. Ratings for the first three operations listed in Table 3 were closely related to those based on light turning cuts. They all had very high coefficients of correlation and determination, and the slopes and intercepts for the mean-square trend lines were similar. The correlation coefficients were similar for end milling and for face milling with turning. Nevertheless, they show that the data fit the trend lines quite well. The coefficients of determination indicate that over 73% of the variances were accounted for. On the other hand, the correlation between machinability ratings for drilling and turning was poor when the 22 types of materials were considered in the statistical analysis. Discrepancies of this type are often caused by differences in the ease with which chips can be removed from the scene of the action. In this case, excluding the data for the five cast irons from consideration resulted in a more meaningful correlation, as indicated in the last line in Table 3 . Table 3 Correlation data for the machinability ratings of different machining operations on ferrous materials Correlation data are based on the machinability ratings(a) of 22 different types of ferrous materials. Correlation statistic of the machinability ratings(a) for the machining operation in the left column and turning with a 1.0 mm (0.04 in.) depth of cut Machining operation compared Trend line with turning at 1.0 mm (0.04 in.) Correlation Coefficient of depth of cut coefficient determination Slope Intercept Turning with a 3.8 mm (0.15 in.) depth of cut

0.9954

0.991

1.02

−2.26

Boring

0.9964

0.993

0.99

−1.13

Reaming

0.9681

0.973

1.08

−10.4

Face milling

0.8827

0.779

0.78

14.9

End milling

0.8583

0.737

0.67

22.5

Drilling(a)

0.6786

0.460

0.61

42.7

Drilling(b) 0.8623 0.740 0.56 37.5 (a) The ratings were calculated from speeds recommended in Ref 1 for machining 22 different types of material with HSS tools. Those speeds were compared with the recommendations for 1212 steel to be processed under similar conditions. The recommended cutting speed for 1212 steel, at 175 HB, was assigned a machinability rating of 100. (b) Ratings from 17 pairs of observations; ratings for cast iron not considered

Venkatesh and Narayanan (Ref 8) used statistical methods to evaluate relationships among machinability ratings determined by drilling, turning, and milling processes. They found better correlations among ratings from different processes than among those


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01 Sep 2005

determined with different tool materials.

Microstructure In general terms, it is possible to differentiate between extremes of metallurgical structure that are easy or difficult to machine. However, when specific cases are encountered, such as differentiating between the machining characteristics of 5130 and 8630 in a particular operation, the presence of a certain microstructure is not, of itself, definitive for selecting one or the other of the two similar steels. This is particularly true when both steels can be heat treated to develop the same type of microstructure. There are rough correlations among hardness, microstructure, and machinability. Experience teaches that, when machining high- or medium-carbon alloy steels such as 4140, the maximum tool life is obtained with workpieces in the annealed condition. Tool wear is accelerated by increases in hardness level. Based on many observations, machinability theory and practice indicate that the optimum conditions or microstructures for machining steels of different carbon contents are usually as follows: Carbon, %

Optimum microstructure

0.06−0.20

As-rolled (most economical)

0.20−0.30

Less than 75 mm (3 in.) in diameter, normalized; 75 mm (3 in.) in diameter and over, as-rolled

0.30−0.40

Annealed to give coarse pearlite, minimum ferrite

0.40−0.60

Coarse lamellar pearlite to coarse spheroidite

0.60−1.00

100% spheroidite, coarse to fine

The above examples have only qualitative utility. An attempt to define the relative machinability of steels by a quantitative measure of microstructure has one additional disadvantage: Even assuming that an optimum structural combination could be predetermined and accurately measured, it would still be necessary to regularly achieve this desired combination on a production basis. This means that all parts would have to be heat treated accurately enough to produce an identical amount and type of microstructure in each. Using commercial steels and production heat treatment practices, such precision is usually impracticable. Among normalized and annealed steels, those with lower hardness and smaller amounts of pearlite can be machined at higher speeds for equal tool lives. By assuming a direct relationship between carbon and pearlite contents. Kronenberg reported that the life of carbide cutting tools decreased as the carbon content of workpiece steels increased (Fig. 5 ). Araki and associates (Ref 9) made a similar analysis on data they obtained on 4135 steel that was heat treated to a variety of structures having different hardnesses. As illustrated in Fig. 6 , the harder specimens caused tool failure in a shorter time than the softer specimens. These results are comparable to those reported by Armarego and Brown, shown in Fig. 7 . An equation of the form: VcHx = C (Eq 5) where H is the Brinell hardness number and x and C are empirical constants, has been used by several investigators. Mayer and Lee (Ref 11) combined Eq 1 and 5 to obtain a relationship between relative tool life and hardness for different steels and tool materials, shown in Fig. 8 . Fig. 5 Effect of carbon and pearlite content on cutting speed. Cutting speed for 60-min tool life in steels containing different amounts of carbon and pearlite; 0.65 mm2 (0.001 in.2) cross-sectional cutting area; carbide tool


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01 Sep 2005

Fig. 6 Effect of hardness on tool life curves. Workpiece: 4135 steel. Tool material: cobalt-tungsten (10% Co, 10% W) high-speed steel per Japanese designation SKH57. Machining conditions: depth of cut = 2.0 mm (0.08 in.); feed rate = 0.2 mm/rev (0.008 in./rev). Source: Ref 9

Fig. 7 Effect of hardness on cutting speed for 30-min tool life, using HSS and carbide tools. Source: Ref 10


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01 Sep 2005

Fig. 8 Effect of hardness on tool life. Relative tool life as a function of relative hardness for three tool materials; the value of x, 1.72, used in constructing these curves is a conservatively estimated maximum. Source: Ref 11

The microconstituents frequently encountered in steels can be identified as contributing to or detracting from the machinability of the steel. Ferrite can be readily cut and causes little tool wear, but it also contributes to the formation of a built-up edge on the tool and a relatively poor surface finish on the workpiece. Spheroidized structures can behave similarly, but large quantities of massive carbide particles can cause significant wear on the tool. Pearlite is harder than ferrite and generally causes greater tool wear; the finer the pearlite plate spacing, the shorter the tool life. A built-up edge is less common when machining pearlite than when machining ferrite. Hard constituents, such as massive carbides or oxides, can be very abrasive to the cutting tool; such particles generally accelerate tool wear. Soft constituents, such as lead or manganese sulfide, generally improve the machinability of the steel. As discussed above, there are correlations between machinability and several interrelated factors, such as composition, microstructure, and hardness; however, it is not clear that there is an exact causal relationship between any one of these factors and machinability. Furthermore, the effects on machining behavior of these factors acting together are generally less than those from differences in tool materials, tool configurations, or the choice of machining process.

Carbon Steels


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01 Sep 2005

Carbon steels nearly always have better machinability than alloy steels of comparable carbon content and hardness. Steels hardened and tempered to hardness levels greater than 300 HB are an exception to this observation; under such conditions, alloy steels have superior machinability, which is usually attributed to, first, the higher tempering temperature required to temper an alloy steel to a specified hardness level and, second, nonuniformity of microstructure due to limited hardenability in carbon steels. Relative machinability ratings for some plain carbon steels are given in Table 4 . Carbon content has a dominant effect on the machinability of carbon steels, chiefly because it governs strength, hardness, and ductility. Increasing the carbon content of steel increases its strength and the unit power consumption for cutting. Data in Fig. 9 show the effects of increasing carbon content and manganese content on unit power consumption. Table 4 Machinability ratings of plain carbon steels Machinability ratings are from the percentage of cutting speed for 1212 steel at a given tool life. SAE/AISI Machinability grade(a) rating

Hardness, HB

1212

100

175

1005

45

95

1006

50

95

1010

55

105

1015

60

111

1017

65

116

1019

70

131

1030

70

149

1038

65

163

1040

60

170

1045

55

179

1045

60(b)

170

1050

45

197

1050

55(b)

189

1065

60(b)

183

1070

55(b)

187

1075

48

192

1085

45(c)

192

1095

45(c)

197

1524

60

163

1536

55

187

1541

45

207

1547

40

207

1547 45(b) 187 (a) Values are for steels cold drawn from the hot-rolled condition, unless otherwise indicated. (b) Annealed, then cold drawn. (c) Spheroidized, then cold drawn

Fig. 9 Effect of carbon content on unit power consumption. Unit power consumption for hot-rolled and cold-drawn steels of two different manganese levels containing various amounts of carbon


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01 Sep 2005

Low-carbon steels containing less than 0.15% C are low in strength in the annealed condition; they machine poorly because they are soft and gummy and adhere to the cutting tools. The machinability of thesegrades can best be improved by work hardening to raise the strength level and lower the ductility. Steels in the 0.15 to 0.03% C range are usually machined satisfactorily in the as-rolled, as-forged, annealed, or normalized condition with a predominantly pearlitic structure. The medium-carbon grades, containing up to about 0.55% C, machine best if an annealing treatment that produces a mixture of lamellar pearlite and spheroidite is utilized. If the structure is not partially spheroidized, the strength and hardness may be too high for optimum machinability. For steels with carbon content higher than about 0.55%, a completely spheroidized structure is preferred. Hardened and tempered structures are generally not desired for machining. Selection of a carbon steel grade within the standard 10xx series is seldom based entirely on machinability, although it may be a factor in selection when other functional requirements can be satisfied by more than one grade. Both tool life and production rate are adversely affected by increases in carbon content. To minimize tool wear and maximize production rate, carbon content should be held to the lowest level consistent with mechanical property requirements. Carbon content also affects surface finish in machining, although its effect can be greatly modified by the nature of the cutting operation or by the cutting conditions. Low values of surface roughness resulting from machining can be most easily achieved with carbon steels containing approximately 0.25 to 0.35% C. The practical significance of carbon content to economy in machining is relatively slight in most selection problems involving hardenable carbon steels because the difference in carbon content between steels with equivalent mechanical properties is unlikely to exceed 0.10%. Therefore, assuming a need for comparable mechanical properties, a choice between 1045 and 1050 would be realistic, while a choice between 1030 and 1090 would not.

Resulfurized Carbon Steels There is a significant improvement in machinability when a resulfurized carbon steel is substituted for a plain carbon steel of approximately the same carbon content. In carbon steels, the sulfur content is ordinarily restricted to a maximum of 0.05%. In the manufacture of resulfurized steels, sulfur is deliberately added to achieve the desired sulfur level. The most common range of sulfur content in resulfurized steels is 0.08 to 0.13%, but some grades permit sulfur content as high as 0.35%. Machinability ratings for standard resulfurized steels are given in Table 5 . Table 5 Machinability ratings of resulfurized and rephosphorized carbon steels, percent of cutting speed for B1112/1212 Machinability rating, %

Hardness, HB

1117

90

137

1118

85

143

1137

70

197

1140

70

170

1141

70

212

1144

80

217

1146

70

187

1151

65

207

1212

100

...

1213

136

...

1215

136

...

12L14

160

163

Grade(a)


Page 915



01 Sep 2005

12L14(b)

190

137

12L14(c)

235

137

12L14(d) 295 137 (a) All values are for cold-drawn steels. (b) Proprietary free-machining variant of 12L14. (c) Proprietary free-machining variant of 12L14 containing bismuth. (d) Proprietary free-machining variants of 12L14 containing bismuth, selenium, or tellurium

Sulfur is added to steel for the sole purpose of decreasing machining costs, either by increasing productivity through greater machining speeds and improved tool life or by eliminating secondary operations through an improvement in finish. Sulfide inclusions, depending on their size, shape, and orientation, improve machining by causing the formation of a broken chip instead of a stringy or continuous chip and by providing a built-in lubricant that prevents the chips from sticking to the tool and undermining the cutting edge. By minimizing this adherence, less power is required, finish is improved, and the speed of machining can often be doubled, compared with machining a similar, nonresulfurized grade. A tightly curled chip that breaks readily is also particularly helpful in milling, deep drilling, tapping, slotting, and reaming because the chip is forced to move within a confined area in these operations. The reduced friction, lower specific power requirements, and improved chip characteristics when machining resulfurized steels all contribute to increased production rates. The advantage of free-machining steels over carbon and alloy grades in terms of unit power consumption is shown in Fig. 10 . The difference is important because almost all the energy of cutting is converted into heat in the cutting zone. Fig. 10 Unit power consumption for free-machining and standard grades of carbon and alloy steels as a function of hardness

The manganese content of resulfurized steels must be high enough to ensure that all the sulfur is present in the form of manganese sulfide (MnS) particles. When a high sulfur content is accompanied by an increase in manganese content, a better surface finish is obtainable, which usually results in an improvement in dimensional accuracy. Control and Effect of Sulfide Morphology. The control of MnS particle shape, size, and distribution is a critical aspect of steelmaking (Ref 12). The particles may remain somewhat globular or may become elongated during rolling to form stringers parallel to the direction of rolling. Figure 11 shows that the size and shape of sulfide particles have pronounced effects on the machinability of steels having similar compositions. The two bars came from different ingots in the same heat. In this case, the difference in machinability was caused by variations in oxygen content and was reflected by the differences in silicon content. The presence of aluminum or other strong deoxidizers changes the shape of sulfide inclusions and may impair machinability. In theory at least, differences in rolling practice that affect the characteristics of sulfide inclusions also influence machinability. Fig. 11 Influence of size and shape of sulfide inclusions on machinability. Two steels, identical in composition except for silicon content, exhibited different machinability ratings that were traced to differences in the size and shape of MnS inclusions. Source: Ref 12


Page 916


Chemical composition, %


01 Sep 2005

Steel A

Steel B

Carbon

0.07

0.07

Manganese

0.94

0.94

Sulfur

0.200

0.200

Phosphorus

0.094

0.093

Silicon

0.009

0.044

176

125

Machinability rating, %

For years, the beneficial effects of large, globular sulfides on machinability have been realized by controlling the silicon contents of semikilled steels (Ref 12). Recently, the effects of inclusion morphology on other characteristics have received considerable attention. Striking benefits resulting from changing inclusions have been reported, but there has been less agreement about the mechanisms through which they act. Royer (Ref 13) reported that a 25% improvement in machinability resulted from increasing the manganese content of type 303 stainless steel. He attributed the improvement to softer, less abrasive sulfide inclusions. Presumably, the higher manganese level may also have affected the total volume and size of the manganese sulfides. Yaguchi (Ref 14) measured thrust and cutting forces in tests on leaded, free-cutting steels. The life to catastrophic failure of the HSS tools improved considerably with inclusion size. The minima in cutting forces shifted toward higher speeds as the sulfides increased in size. Yaguchi concluded that the size of the inclusions influenced the temperature distribution at the tip of the tool and indirectly affected the formation of the built-up edge on the tool. It is widely known that the presence or absence of a built-up edge influences tool life. Abeyama and associates (Ref 15) controlled the morphology of sulfides in resulfurized free-cutting steels through additions of tellurium. Treatments producing Te/S ratios of 0.2 were most effective in improving machinability, and these treatments also reduced the anisotropy of mechanical properties. The researchers attributed both effects to the influence of tellurium on the melting or softening temperatures of the inclusions. Raising the melting point of the sulfides by alloying would minimize the elongation of the mixed sulfides during hot rolling. Katayama (Ref 16) found that sulfide shape also affected the machinability of continuously cast billets. This subject is of interest because steels made by that practice, rather than from ingots, have smaller as-cast inclusions that are not elongated as much during hot rolling because the reductions are smaller. However, this difference between ingots and continuously cast steel may not always manifest itself in terms of machinability. Welburn and Naylor (Ref 17) reported that the machinability of 1144 steel made from continuously cast billets is equivalent to that of bars made from ingots. Their conclusions were based on both laboratory tests and experience by customers. Economic factors influence the use of resulfurized steels because resulfurized grades cost more than plain carbon steels. For example, 25 mm (1 in.) round hot-rolled bars of 1117 steel cost approximately 8% more than plain carbon steel bars of a composition that is similar except for the sulfur content. Nevertheless, the economic benefits of machining resulfurized steels are large enough to justify annual purchases of 2 million tons of free-cutting steel in the United States. The use of resulfurized grades depends on whether the higher steel cost can be offset by lower machining cost. Considering only the cost of removing chips, the use of free-cutting steels can seldom be justified if less than about 10% of the bar is removed in machining the parts. However, as the amount of metal machined off approaches or exceeds 20%, the resulfurized grades should be considered. When the required finish can be obtained in a primary operation with a free-cutting grade but a secondary operation would be necessary for a grade that is not free cutting, the 11xx series can be justified for a smaller amount of metal removal. In such parts, the savings in handling and the elimination of an operation usually more than offset the extra cost of the


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01 Sep 2005

steel without consideration of machining speed or tool life. Figure 12 illustrates how the substitution of heavily drawn and stress-relieved 1144 for heat-treated 1046 can significantly increase the number of parts machined between regrinds of the tools and therefore the cost of manufacturing the part. Fig. 12 Comparison between machining 1046 and 1144

Carbon Steels With Other Additives Phosphorus, as well as sulfur, is often added to improve the machining characteristics of low-carbon steels. The phosphorus range for 12L14 and 1215 is 0.04 to 0.09%. The phosphorus limits are 0.07 to 0.12% for other steels in the 12xx series. The limits are set because phosphorus, like carbon, increases the hardness and strength of the steel. Consequently, excessive phosphorus contents impair machining characteristics and some other properties of steel. Phosphorus is soluble in iron and increases the strength of ferrite, an effect that promotes chip breaking in cutting operations. The phosphorus helps to avoid the formation of long, stringy chips in some operations and may result in a better surface finish. Nitrogen. The effects of nitrogen on the machining characteristics of 12L14 steel were studied by Watson and Davies (Ref 18). They found that nitrogen adversely affected the life of HSS tools used for turning and form cutting. This effect was attributed to strain age hardening, caused by nitrogen, before testing. No beneficial effects of nitrogen were detected. Selenium and tellurium additions improve machinability but are not available in standard grades of steel. These additions are expensive (selenium treatment increases the cost of steel by about 15%). When they are used, they are often used in combination with sulfur or lead. Typical percentages of either element would be 0.04 or 0.05%. Both elements seem to exert beneficial effects by promoting the retention of globular-shaped sulfide-type inclusions. For the same reason, they are considered to have a less deleterious effect than sulfur on mechanical properties. The data in Fig. 13 show that the effect of tellurium on machinability can be appreciable. The data were obtained on steels with a nominal tensile strength of 1035 MPa (150 ksi) by turning with a form tool and measuring the diameters of successive parts. The presence of 0.042% Te quadrupled the number of parts made between tool changes and improved the surface finish. Tata and Sampsell (Ref 19) reported that selenium is even more effective than tellurium in improving the machinability of steels, particularly alloy steels. Fig. 13 Effect of tellurium on tool wear. Tool wear, as measured by part growth, in multiple-operation machined parts of quenched and tempered 4142 and a similar grade with tellurium. Cutting speed was 0.5 m/s (99 sfm). Source: Ref 11

Calcium additions improve the machining characteristics of steels fully deoxidized with aluminum. The cost of the special


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01 Sep 2005

treatment is relatively modest. Steels made by aluminum deoxidation practices ordinarily contain small inclusions of aluminum silicate in quantities essentially independent of the amount of aluminum added to the steel. The inclusions are often assumed to be alumina, and the poorer machinability of aluminum-killed steels, compared to steels deoxidized with silicon, is often attributed to the supposedly abrasive effects of the inclusions. The validity of this explanation is debatable. Calcium additions result in larger inclusions consisting of calcium-aluminum silicates. Joseph and Tipnis (Ref 20) considered these inclusions to be softer and less abrasive than those in steels not treated with calcium and concluded that such attributes benefited machinability. Tests on a series of 1045 steels led Subramanian and Kay (Ref 21) to the same conclusions. Abeyama and colleagues (Ref 15, 22) demonstrated that calcium treatments benefit the machining characteristics of several types of steel. Their studies on heavy-duty steels led them to believe that machinability is affected by the chemical composition of the inclusions and by temperatures at the tool point. Fombarlet (Ref 23) attributed the better machining properties of 1048, 4142, and 8620 steels to appropriate treatments with calcium.

Leaded Carbon and Resulfurized Steels The addition of lead to carbon steels is another means of increasing the machinability of the steels and improving the surface finish of machined parts. Lead is added to the molten steel during teeming of ingots or, sometimes, to the ladle. Because lead is insoluble, or nearly so, in molten steel, a fine dispersion of lead particles develops as the steel solidifies. The lead is usually found near or surrounding the sulfide inclusions. On special order, nearly all carbon steels in the 10xx and 11xx series can be produced with 0.15 to 0.35% Pb. The grades are identified by inserting the letter "L" between the second and third numerals of the grade designation, for example, 10L45. It is generally believed that lead has a minimal effect on the yield or ultimate strength, ductility, or fatigue properties of steels at room temperature and moderate strength levels. Lead can also be added to alloy steels to improve machinability without sacrificing room-temperature mechanical properties. Environmental considerations may restrict the manufacture or use of leaded steels. Leaded steels cost about 5% more than similar nonleaded compositions. Because requirements of machinability and finish are the only reasons for using leaded grades, the extra cost for these steels must be justified by either or both of these factors. As with resulfurized grades of carbon steel, consideration must be given to the amount of stock being removed. For example, if only 10% of the bar weight is to be removed in machining, the extra cost for leaded grades may not be justifiable. However, if 20% or more of the bar is converted to chips, the leaded grades should be considered. The problem of surface finish requires consideration of specific parts, and no general statements are valid. In a case study of one particular part, the finish obtained in drilling and reaming standard 1050 steel was marginal, and parts were frequently rejected and subsequently reworked. The use of leaded 1050 corrected this condition, and the added cost of the steel was justified by finish alone. Most of the resulfurized grades can be produced with an addition of 0.15 to 0.35% Pb. The lead addition augments the effect of sulfur, permitting a further increase of machining speed and better finish. For screw machine parts where more than 50% of the bars become chips, there may be justification for the higher cost, especially where there are high finish requirements. The increased cost of leaded steel is unjustified when machine tools are already being operated at maximum speed on plain resulfurized steels. On the other hand, machine tools that have been designed for higher speeds can take advantage of the leaded resulfurized grades. In another case study, surface finish was the deciding factor in selecting a leaded steel for parts about 13 mm (1=2in.) in diameter and 75 mm (3 in.) long made from 1141 steel on a multiple-operation machine. Finish requirements on the ends of the parts could not be met in the cutoff operation, and a secondary facing operation was necessary. A change to leaded 1141 steel provided just enough better finish from the cutoff operation so that the finish requirement could be met, and the elimination of the facing operation justified the cost of the leaded grade. The example shown in Fig. 14 illustrates the importance of machinability of steel in determining the cost of small parts. A change from 1213 to 12L13 nearly doubled the production rate of these bushings and reduced the total cost per part, even though the 12L13 steel was more costly. Fig. 14 Effect of lead content on multiple-operation machining. Graphs show the production rate and relative cost of machining bushings from 1113 and 12L13 steel with cutting speeds chosen for an 8-h tool life.


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Reh and coworkers (Ref 24) described the characteristics of lead-free, high-sulfur steels containing bismuth. They found bismuth to be more efficient than lead in reducing cutting forces and in improving surface finish. They attributed the pronounced improvements caused by bismuth to the spheroidization of inclusions and to the liquid-metal embrittlement of grain boundaries, by liquid bismuth, during cutting.

Carburizing Steels It is difficult to evaluate the relative machining economy of plain carbon carburizing steels, except in terms of specific parts, largely because these steels posses conflicting properties that may either promote or detract from economy in machining. On the one hand, their low carbon content may be beneficial to tool life and production rate. However, their relatively soft, gummy structure results in a tearing action in cutting, which is harmful to surface finish and dimensional accuracy. It is this conflict in properties that sometimes makes it advantageous to machine or partially machine these steels in the carburized condition. Although higher carbon content adversely affects tool life, the higher carbon areas are more controllable in terms of surface finish and dimensional accuracy. In operations such as gear cutting, the finish and accuracy may be more decisive factors than tool wear or production rate.

Through-Hardening Alloy Steels The steels chosen for parts that must be hardened throughout must contain enough carbon to achieve the desired hardness after quenching and sufficient alloy content to obtain the desired percentage martensite in the thickest section of the part. The combination of carbon and alloy contents can make these steels difficult to machine. The extent of the difficulties encountered in machining these steels depends primarily on the microstructure and hardness of the steel and secondarily on its alloy content. Figure 15 illustrates the magnitude of the variations in tool life that may be expected from differences in hardness and microstructure in a single steel, 4340. A similar effect for 4135 steel is shown in Fig. 6 . Machinability ratings for several alloy steels are given in Table 6 . Table 6 Machinability ratings for alloy steels compared to 1212 steel The machinability rating of 1212 steel is assigned at 100. Machinability rating(a)

Hardness, HB

1330

55(b)

179−235

1340

50(b)

183−241

1345

45(c)

183−241

4024

75(c)

156−207

4028

75(c)

167−212

4042

65(b)

179−229

4130

70(b)

187−229

4140

65(b)

187−229

41L40

85(b)

185−230

4150

55(b)

187−240

4340

50(b)

187−240

4620

65(c)

183−229

Grade


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50B40

65(b)

174−223

50B60

55(d)

170−212

5130

70(b)

174−212

5140

65(b)

179−217

5160

55(d)

179−217

51B60

55(d)

179−217

50100

40(d)

183−240

51100

40(d)

183−240

52100

40(d)

183−240

8115

65(c)

163−202

81B45

65(b)

179−223

8630

70(b)

179−229

8620

65(c)

179−235

86L20

85(c)

...

8660

55(d)

179−217

8645

65(b)

184−217

86B45

65(b)

184−217

8740 65(b) 184−217 (a) Ratings are for cold-finished bars. (b) Microstructure composed of ferrite and lamellar pearlite. (c) Microstructure composed mainly of acicular pearlite and bainite. (d) Microstructure composed primarily of spheroidite

Fig. 15 Effect of hardness on tool life curves. Workpiece: 4340 steel. Tool material: C6 carbide. Source: Ref 25

Hardness,


Heat

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Sample

HB

½A

206

01 Sep 2005

treatment Spheroidized

&fcirc; B

221

Annealed

∆C

321

Normalized

´D

400

Hardened and tempered

²E

500


³F

515


A comparison of the machinability ratings with the compositions of these steels indicates that all of the alloying elements that increase the hardenability of the steel decrease machinability; ferrite-strengthening elements such as nickel and silicon decrease the machinability more than equivalent amounts of carbide-forming elements such as chromium and molybdenum. It is not uncommon for heat-treating considerations to overshadow both machining and material costs in the selection of steel. On occasion, heat-treating responses may dictate the selection of a less machinable or a more expensive steel so that the lowest total costs can be realized. The sulfur content of through-hardening alloy steels can significantly affect machining behavior. Variations in residual sulfur level can account for unexplained differences in the machining behavior of different lots of the same material. Many grades of hardenable alloy steels can be obtained in the resulfurized condition. The differences in tool life and cutting speed between standard and high-sulfur 4150 steels are substantial. Tests by Field and Zlatin (Ref 26) showed that raising the sulfur content from 0.04 to 0.09% increased the cutting speed for 60-min tool life by 25%. Alloy steels containing lead are available and useful. As indicated in Table 6 , the machinability rating of the leaded grade 41L40 is 85, while the rating for 4140 is only 65. The performance of these two grades in several machining operations is indicated in Table 7 . The data are from a case study described in Ref 26. Table 7 Effect of lead on cutting speed and tool life in machining alloy steels Standard 4140

Leaded 4140

Hardness, HB

300

300

Cutting speed, rev/min

321

495

Operation Turning

Feed, mm/rev (in./rev) Tool life, parts per tool grind

0.30

0.30

(0.012)

(0.012)

4

18−20

460

740

Turning Cutting speed, rev/min Feed, mm/rev (in./rev)

0.15

0.23

(0.006)

(0.009)

8.76

10.55

(28.75)

(34.6)

0.10

0.15

Drilling(a) Cutting speed, m/min (sfm) Feed, mm/rev (in./rev)

(0.004) (0.006) (a) In drilling standard 4140 steel, the 19 mm (3=4in.) diam hole jammed with chips and the drill had to be removed frequently for cleaning. When using leaded 4140 steel, the entire depth was drilled without removing the tool.

Another important factor that can affect the choice of steel for a through-hardening application is the effect of alloying elements added for machinability on the mechanical properties of the steel. These steels are often used at high-strength levels, where the deleterious effects of inclusions, particularly on transverse properties, might not be permissible. The effect of sulfur, in the amounts usually specified for enhanced machinability, is generally considered to be more damaging than that of lead. For some applications, neither machinability additive can be tolerated.

Cold-Drawn Steel Cold drawing generally improves the machinability of steels containing less than about 0.2% C. The improvement is most noticeable in plain carbon steels, as shown in Fig. 16 . The machinability of higher-carbon steels, or alloy steels, is less affected by cold work. This improvement in machinability may be attributed to reduced cutting forces and/or the characteristics of chip removal. Kopalinsky and Oxley (Ref 28) found that cold drawing lowered the cutting forces and improved the tool life and surface finish of low-carbon steels. Screw machine tests by Yaguchi (Ref 29) showed that the workpiece surface finish improved


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continuously with increases in reduction in area up to 29%. These effects were not characteristic of steels with high nitrogen contents (Ref 18). The improved machinability of cold-drawn steels can also be attributed to the decrease in ductility that results from cold working; thus, the chips are generally not long and stringy. Fig. 16 Effect of cold drawing on tool life. Workpiece: 1016 steel, 25 mm (1 in.) in diameter. Machining conditions: multiple-operation machined with a cutting speed of 0.73 m/s (144 sfm). Source:Ref 27

Cold-finished bars have closer dimensional tolerances, better surfaces, and usually, higher strength than hot-finished bars. The first two factors may be significant in the selection of steels to be machined in multiple-operation machines or other high-production equipment. These considerations are discussed in the article "Cold-Finished Steel Bars" in this Volume. The machining characteristics of cold-drawn steels are only rarely a decisive criterion for selection. The extra strength obtained with cold-drawn steel may be more important from a cost standpoint, because it is often high enough to eliminate the need for heat treatment. REFERENCES 1. 2. 3. 4.

Machining Data Handbook, 3rd ed., Metcut Research Associates Inc., 1980 "Life Tests for Single-Point Tools of Sintered Carbide," B94.36-1956 (R 1971), American National Standards Institute "Tool Life Testing With Single-Point Turning Tools," ANSI/ASME B94.55M-1985, American National Standards Institute J.F. Kahles, Elements of the Machining Process, in Metals Handbook: Desk Edition, American Society for Metals, 1985, p 27.10 5. "Machining Performance of Ferrous Metals Using an Automatic Screw/Bar Machine," E 618-81-03.01, Annual Book of ASTM Standards, American Society for Testing and Materials 6. F.W. Boulger, Influence of Metallurgical Properties on Metal-Cutting Operations, Society of Manufacturing Engineers, 1958 7. F.W. Boulger and H.J. Grover, Machinability Can Be Related to Composition, Tool Eng., Vol 40, March 1958 8. V.C. Venkatesh and V. Narayanan, Machinability Correlations Among Turning, Milling and Drilling Processes, Ann. CIRP, Vol 35 (No. 1), 1986, p 59−62 9. T. Araki et al., Some Results of Cooperative Research on the Effect of Heat Treated Structure on the Machinability of a Low Alloy Steel in Influence of Metallurgy on Machinability, V.A. Tipnis, Ed., American Society for Metals, 1975, p 381−395 10. E.J.A. Armarego and R.H. Brown, The Machining of Metals, Prentice-Hall, 1969 11. J.E. Mayer, Jr., and D.G. Lee, Influence of Machinability on Productivity and Machining Cost, in Influence of Metallurgy on Machinability, V.A. Tipnis, Ed., American Society for Metals, 1975, p 31−54 12. F.W. Boulger et al., Superior Machinability of MX Steel Explained, Iron Age, Vol 167, 17 May 1951, p 90−95 13. W.E. Royer, Making Stainless More Machinable⎯303 Super X, Autom. Mach., Vol 47 (No. 5), May 1986, p 47−49 14. H. Yaguchi, Effect of MnS Inclusion Size on Machinability of Low-Carbon, Leaded, Resulfurized Free-Machining Steel, J. Appl. Metalwork., Vol 3 (No. 3), July 1986, p 214−225 15. S. Abeyama et al., Development of Free Machining Steel With Controlled-Shape Sulfides, Bull. Jpn. Inst. Met., Vol 24 (No. 6), 1985, p 518−520 16. S. Katayama et al., Improvements in Machinability of Continuously-Cast, Low-Carbon, Free-Cutting Steels, Trans. ISI, Vol 25 (No. 9), Sept 1985, p B229 17. R.M. Welburn and D.J. Naylor, Production and Machinability of Billet-Cast Medium Carbon High Sulfur (Over 0.08%) Free-Machining Steels, in Proceedings of the Conference on Continuous Casting, Institute of Metals, 1985 18. J.D. Watson and R.H. Davies, The Effects of Nitrogen on the Machinability of Low-Carbon Free-Machining Steels, J. Appl. Metalwork., Vol 3 (No. 2), 1984, p 110−119 19. H.J. Tata and R.E. Sampsell, Effects of Additions on Machinability and Properties of Alloy-Steels Bars, Paper 730114, Trans. SAE, Vol 82, 1973 20. R.A. Joseph and V.A. Tipnis, The Influence of Non-Metallic Inclusions on the Machinability of Free-Machining Steels, in Influence of Metallurgy on Machinability, V.A. Tipnis Ed., American Society for Metals, 1975, p 55−72 21. S.V. Subramanian and D.A.R. Kay, Inclusions and Matrix Effects on the Machinability of Medium Carbon Steels, in


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Conference Proceedings, Ottawa, Ontario, Canada, Canadian Government Publishing Centre, 1985 22. T. Kato, S. Abeyama, A. Kimura, and S. Nakamura, The Effect of Ca Oxide Inclusions on the Machinability of Heavy Duty Steels, in The Machinability of Engineering Materials, R.W. Thompson, Ed., Conference Proceedings, 13−15 Sept (Rosemont, IL), American Society for Metals, 1983, p 323−337 23. J. Fombarlet, Improvement in the Machinability of Engineering Steels Through Modification of Oxide Inclusions, in The Machinability of Engineering Materials, 13−15 Sept (Rosemont, IL), R.W. Thompson, Ed., Conference Proceedings, American Society for Metals, 1983, p 366−382 24. B Reh, U. Finger et al., Development of Bismuth-Alloyed High Performance Easy Machining Steel, Neue Hütte, Vol 31 (No. 9), Sept 1986, p 327−330 25. N. Zlatin and J. Christopher, Machining Characteristics of Difficult to Machine Materials, in Influence of Metallurgy on Machinability, V.A. Tipnis, Ed., American Society for Metals, 1975, p 296−307 26. M. Field and N. Zlatin, Evaluation of Machinability of Rolled Steels, Forgings and Cast Irons, Machining⎯Theory and Practice, American Society for Metals, 1950, p 341−376 27. J.D. Armour, Metallurgy and Machinability of Steels, Machining⎯Theory and Practice, American Society for Metals, 1950, p 123−168 28. E.M. Kopalinsky and P.L.B. Oxley, Predicting Effects of Cold Working on Machining Characteristics of Low-Carbon Steels, J. Eng. Ind. (Trans. ASME), Vol 109 (No. 3), 1987, p 257−264 29. H. Yaguchi and N. Onodera, Effect of Cold Working on the Machinability of AISI 12L14 Steel, in Strategies for Automation of Machining: Materials and Processes, Proceedings of an International Conference (Orlando, FL), ASM INTERNATIONAL, 1987, p 15−26


Page 924


Weldability of Steels

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Fabrication Characteristics of Carbon and Low-Alloy Steels Weldability of Steels S. Liu, Center for Welding and Joining Research, Colorado School of Mines; J.E. Indacochea, Department of Civil Engineering, Mechanics, and Metallurgy, University of Illinois at Chicago THE MAIN OBJECTIVE of this article is to survey the factors controlling the weldability of carbon and low-alloy steels in arc welding. A good understanding of the chemical and physical phenomena that occur in the weldment is necessary for welding modern steels. Therefore, the influence of operational parameters, thermal cycles, and metallurgical factors on weld metal transformations and the susceptibility to hot and cold cracking are discussed. Common tests to determine steel weldability are also described. The carbon and low-alloy steels group comprises a large number of steels that differ in chemical composition, strength, heat treatment, corrosion resistance, and weldability. These steels can be further divided into subgroups: • • • • •

Carbon steels High-strength low-alloy (HSLA) steels Quenched and tempered (QT) steels Heat-treatable low-alloy (HTLA) steels Precoated steels

This article addresses only the basic principles that affect the weldability of carbon and low-alloy steels. More detailed information concerning the other aspects of welding, such as joint design, defects, and failure in weldments and the influence of these factors on different groups of steels, can be found in Volumes 1and 6 of the 9th Edition Metals Handbook and Volume 11 of ASM Handbook, formerly 9th Edition Metals Handbook and in the "Selected References" at the end of this article.

Characteristic Features of Welds Single-Pass Weldments. To understand weldability, it is necessary to recognize the various weld regions. In the case of a single-pass bead, the weldment is generally divided into two main regions: the fusion zone, or weld metal, and the heat-affected zone (HAZ), as shown in Fig. 1 . Within the fusion zone, the peak temperature exceeds the melting point of the base metal, and the chemical composition of the weld metal will depend on the choice of welding consumables, the base metal dilution ration, and the operating conditions. Fig. 1 Various regions of a bead-on-plate weld

Under conditions of rapid cooling and solidification in the weld metal, alloying and impurity elements segregate extensively to the center of the interdendritic or intercellular regions and to the center parts of the weld, resulting in significant local chemical inhomogeneities. Accordingly, the transformation behavior of the weld metal may be quite different from that of the base metal,


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even when the bulk chemical composition is not significantly changed by the welding process. The typical anisotropic nature of the solidified weld and structure is also shown in Fig. 1 . The chemical composition remains largely unchanged in the HAZ because the peak temperature remains below the melting point of the parent plate. Nevertheless, considerable microstructural change takes place within the HAZ during welding as a result of the extremely harsh thermal cycles. The material immediately adjacent to the fusion zone is heated high into the austenitic temperature range. The microalloy precipitates that development in the previous stages of processing will generally dissolve, and unpinning of austenite grain boundaries occurs with substantial growth of the grains, forming the coarse-grain HAZ. The average size of the austenite grains, which is a function of the peak temperature attained, decrease with increasing distance from the fusion zone. The cooling rate also varies from point to point in the HAZ; it increases with increasing peak temperature at constant heat input and decreases with increasing heat input at constant peak temperature. Because of varying thermal conditions as a function of distance from the fusion line, the HAZ is actually composed of coarse-grain zones (CGHAZ), fine-grain zones (FGHAZ), intercritical zones (ICHAZ), and subcritical zones (SCHAZ). The various HAZ regions of a single-pass low-carbon steel butt weld are shown in Fig. 2 . Fig. 2 Various regions of the HAZ of a single-pass low-carbon steel weld metal with 0.15 wt% C

In multipass weldments, the situation is much more complex because of the presence of reheated zones within the fusion zone, as shown in Fig. 3 . The partial refinement of the microstructure by subsequent weld passes increases the inhomogeneity of the various regions with respect to microstructure and mechanical properties. Reaustenitization and subcritical heating can have a profound effect on the subsequent structures and properties of the HAZ. Toughness property deterioration is related to small regions of limited ductility and low cleavage resistance within the CGHAZ that are known as the localized brittle zones (LBZ). Localized brittle zones consist of unaltered CGHAZ, intercritically reheated coarse-grain (IRCG) heat-affected zone, and subcritically reheated coarse-grain (SRCG) heat-affected zone. At an adjacent fusion line, that LBZs may be aligned, as shown in Fig. 3 . The aligned LBZs offer short and easy paths for crack propagation. Fracture occurs along the fusion line. Fig. 3 Overlapping of HAZ to form localized brittle zones aligned along the fusion line


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Metallurgical Factors That Affect Weldability Hardenability and Weldability. Hardenability in steels is generally used to indicate austenite stability with alloy additions. However, it has also been used as an indicator of weldability and as a guide for selection a material and welding process to avoid excessive hardness and cracking in the HAZ. Steels with high hardness often contain a high volume fraction of martensite, which is extremely susceptible to cracking during processing. Hardenability is also used to indicate the susceptibility of a steel to hydrogen-induced cracking. Traditionally, empirical equations have been developed experimentally to express weldability. Carbon equivalent (CE) is one such expression; it was developed to estimate the cracking susceptibility of a steel during welding and to determine whether the steel needs pre- and postweld heat treatment to avoid cracking. Carbon equivalent equations do include the hardenability effect of the alloying elements by expressing the chemical composition of the steel as a sum of weighted alloy contents. To date, several CE expressions with different coefficients for the alloying elements have been reported. The International Institute of Welding (IIW) carbon equivalent equation is:

(Eq 1) where the concentration of the alloying elements is given in weight percent. It can be seen in Eq 1 that carbon is the element that most affects weldability. Together with other chemical elements, carbon may affect the solidification temperature range, hot tear susceptibility, hardenability, and cold-cracking behavior of a steel weldment. Figure 4 summarizes the CE and weldability description of some steel families. Because of the simplification and generalization involved in Fig. 4 , it should be used


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cautiously for actual welding situations. Fig. 4 Weldability of several families of steels as a function of carbon equivalent. 1, Mo; 2, Cr + Ni + Mo + Si, and so on; 3, Cr or V or Ni + Si, and so on

The application of CE expressions is also empirical. For example, the IIW carbon equivalent equation has been used successfully with traditional medium-carbon low-alloy steels. Steels with lower CE values generally exhibit good weldability. When the CE of a steel is less than 0.45 wt%, weld cracking is unlikely, and no heat treatment is required. When the CE is between 0.45 and 0.60 wt%, weld cracking is likely, and preheat in the range of approximately 95 to 400 °C, (200 to 750 °F) is generally recommended. When the CE of a steel is greater than 0.60 wt%, there is a high probability that the weld will crack and that both preheat and postweld heat treatments will be required to obtain a sound weld. However, Eq 1 does not accurately correlate with the microstructures and properties of newly developed low-carbon microalloyed steels over extended alloy ranges. Thus, new expressions based on solution thermodynamics and kinetic considerations were developed to obtain better predictions of the alloy behavior and weldability of low-carbon low-alloy steels. Complex interactive terms, rather than simple additive forms, are included in these equations. An example of one such expression is: CE = k1C[1 + k2C + k3Mn + ... Ã+ k11 ln C + k22C ln C + k33Mn ln Mn Ã+ ... + k111CMn + ...] (Eq 2) where k1, k2, ..., and so on, are the weighted coefficients multiplied to the concentration of the alloying elements. Nonlinear terms such as ln Xi, Xi ln Xi, and XiXj represent the interaction effect among the alloying elements Xi and Xj. Equations with these nonlinear terms are more useful in predicting arc welding behavior. Several expressions are also available for other steel groups with a wider range of alloying elements and with different prior heat treatments, hydrogen contents, and weld hardnesses. Recently, expressions that include fabrication conditions such as heat input, cooling rate, joint design and restraint conditions have also been proposed. An example of this type of equation is:

(Eq 3) where PH is the cracking susceptibility parameter, H is the concentration of hydrogen (in parts per million), Rf is the restraint stress (in megapascals), and:


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01 Sep 2005

(Eq 4) The thickness of the part being welded can also be related to CE as a compensated carbon equivalent (CCE) as follows: CCE = CE + 0.00254e (Eq 5) where e is the thickness of the part (in millimeters). Equations 3 , 4 , and 5 are valid only for specific ranges of chemical composition and welding conditions. Nevertheless, despite the different forms and terms included in the predictive equations, the main objective remains that of estimating the weldability and cracking susceptibility of the material. Weld Metal Microstructure. Inherent in the welding process is the formation of a pool of molten metal directly below a moving heat source. The shape of this molten pool is determined by the flow of both heat and metal, with melting occurring ahead of the heat source and solidification occurring behind it. Heat input determines the volume of molten metal and therefore the dilution and weld metal composition, as well as the thermal conditions under which solidification takes place. Also important to solidification is the crystalline growth rate, which is geometrically related to weld travel speed and weld pool shape. Thus, weld pool shape, weld metal composition, cooling rate, and growth rate are all factors that are interrelated with heat input, which in turn will affect the solidification microstructure and the tolerance of the weldment to hot cracking. Incipient melting at base metal grain boundaries immediately adjacent to the fusion zone allows these grains to serve as seed crystals for epitaxial grain growth during weld metal solidification. The continuous growth of the epitaxial grains results in large columnar grains whose boundaries provide easy paths for crack propagation. An elongated weld pool will yield straight and broad columnar grains, which promote the formation of centerline cracking because of impurity segregation, mechanical entrapment of inclusions, and the shrinkage stresses that develop during solidification. Epitaxial columnar growth is particularly deleterious in multipass welds where grains can extend continuously from one weld bead to another. Hot tears originate near the liquid/solid interface when strains from solidification shrinkage and thermal contraction cause rupture of the liquid films of low melting point located at grain boundaries. The susceptibility of an alloy to hot tearing is related to its inability to accommodate strain through dendrite interlocking as well as the tendency of tears to backfill with the remaining liquid. The time interval during which liquid films can exist in relation to the rate of strain generation may also play a role in hot tear susceptibility. Ferrous alloys can be hot tear sensitive depending on the amount of phosphorus and sulfur impurities they contain. Carbon and nickel are also known to influence hot cracking in steel welding. When the solidified steel weld metal cools down, solid-state transformation reactions may occur. As in solidification, the two main factors that determine the final microstructure are the chemical composition and thermal cycle of the weld metal. In most structural steels, weld metal will solidify as δ-ferrite. At the peritectic temperature, austenite will form from the reaction between liquid weld metal and δ-ferrite, and subsequent cooling will lead to the formation of α-ferrite. During the austenite-to-ferrite transformation, proeutectoid ferrite forms first along the austenite grain boundaries; this is known as grain-boundary ferrite. Subsequent to grain-boundary ferrite formation, ferrite sideplates develop in the form of long needlelike ferrite laths that protrude from the allotriomorphs. A coarse austenite grain size and a low carbon content, in combination with a relatively high degree of supercooling, are found to promote ferrite sideplate formation. These laths can be properly characterized by their length-to-width aspect ratios; values above 10:1 are common. As the temperature continues to drop, intragranular acicular ferrite will nucleate and grow in the form of short laths separated by high angle boundaries. The inclination between orientations of adjacent acicular ferrite laths is usually larger than 20°. The random orientation of these laths provides good resistance to crack propagation. Acicular ferrite laths have aspect ratios ranging from 3:1 to 10:1. During proeutectoid ferrite formation, carbon is rejected continuously from the ferrite phase, enriching the remaining austenite, which later transforms into a variety of constituents, such as martensite (both lath and twinned), bainite, pearlite, and retained austenite. Because of the acicular nature of the bainite laths, they can also be described by their aspect ratio, with values similar to those of Widmanstätten side-plates. More frequently, however, bainite laths occur in the form of packets associated with grain boundaries. Figure 5 illustrates the microstructure of a low-carbon steel weld metal. Fig. 5 Weld metal microstructure of HSLA steel. A, grain-boundary ferrite; B, acicular ferrite; C, bainite; D, sideplate ferrite


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Heat-Affected Zone Microstructure. In terms of microstructure, long bainite laths with alternate layers of connected martensite islands are generally found in the CGHAZ of high-strength low-alloy steel weldments. Martensite islands (martensite-retained austenite constituents) are formed because of the enrichment of carbon in austenite in the intercritical zone. Coarse austenite grain size in the near-fusion region of the HAZ can suppress high-temperature transformation products in favor of martensite and bainite upon cooling. Upper bainite has a relatively high transformation temperature and is stable relative to the thermal cycles subsequent to those of the first pass. Fluctuation of the chemical composition of the microalloying elements could also contribute to carbon equivalent change and to the amount of hard martensite present in the CGHAZ. In the FGHAZ, even though the peak temperature attained is above thermal cycle Ac3, it is still well below the grain-coarsening temperature. The smaller prior-austenite grain size and subsequent ferrite transformation produce a refined microstructure having grains smaller than those of the parent material. The microstructure is similar to that of a normalized steel, with considerable toughness. Only partial transformation takes place in the ICHAZ, resulting in a mixture of austenite and ferrite at the peak temperature of the thermal cycle. Upon cooling, the austenite in a matrix of soft ferrite decomposes, and the final microstructure depends on the bulk and local composition of the alloying elements. The cooling rate is also an important factor in determining the amount of martensite and bainite in the ferrite matrix. In the SCHAZ, no observable microstructural changes are observed. Some spheroidization of carbides may occur. Upon reheating by subsequent weld passes, precipitates or preprecipitate clusters may form, reducing the toughness. Irregularly shaped particles may also coalesce and strain the surrounding matrix, further lowering the toughness. During HAZ thermal cycles between Ac1 and Ac3, the austenite becomes enriched with carbon, which, upon cooling, transforms to martensite islands. In the as-welded condition, this transformation affects the IRCG region more than the other reheated zone. Figure 6 illustrates the different phases that can be found in a low-carbon steel HAZ.


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Fig. 6 Heat-affected zone exhibiting a wide variety of microstructures in the intercritical and subcritical regions. A, spheroidized carbides; B, bainite and martensite

Chemical Composition Effect. The presence of a certain phase in the final microstructure of a weldment can be explained by means of a continuous cooling transformation (CCT) diagram, which is formed by two sets of curves: the percent transformation curves and the cooling curves. The percent transformation curves define the regions of stability of the different phases. The cooling curves represent the actual thermal conditions that the weld experienced. The intersection of these two sets of curves determines the final microstructure of the different weld zones. Figure 7 illustrates the use of a CCT diagram to determine the microstructure of a low-carbon low-alloy steel weld metal. Fig. 7 Continuous cooling transformation diagram for an HSLA steel weld metal showing the effect of cooling rate and chemical composition on microstructure. CR, cooling rate


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Hardenability elements, such as carbon, manganese, chromium, and molybdenum, suppress the start of austenite decomposition to lower temperatures. This is equivalent to pushing the transformation curves to the right side of the CCT diagram, resulting in a refined microstructure. Inclusion formers, such as oxygen and sulfur, accelerate the austenite-to-ferrite transformation by providing more nucleation sites for the reaction to initiate at higher temperatures. Faster cooling has the same effect as an increase in hardenability elements, while a slower cooling rate acts in the same direction as a decrease in hardenability agents or an increase in nucleation site providers. Because the cooling rate varies from point to point in the HAZ, the microstructure also changes accordingly, with martensite and bainite in regions close to the fusion line. Preweld and Postweld Heat Treatments. In the welding of carbon and low-alloy steels, the final microstructure of the weldment is primarily determined by the cooling rate from the peak temperature. Because the alloy level in carbon and low-alloy steel is low, the major physical properties of the steel are not affected. Thus, temperature gradient and heat input become the important parameters in weld metal microstructural evolution. A slower cooling rate decreases shrinkage stress, prevents excessive hardening, and allows time for hydrogen diffusion. Cooling rate (CR) is of particular importance and is a function of the difference in temperature, ∆T, as well as the thermal conductivity of the material, k. The cooling rate can be expressed for thinplate and thick-plate welding, respectively, as: CR ∝ k ∆T3 CR ∝ k ∆T2 (Eq 6) During preheating, the initial temperature of the plate increases, decreasing the cooling rate and the amount of the hard phases, such as martensite and bainite, in the weld microstructure. For the welding of hardenable steels, it is important to determine the critical cooling rate (CCR) that the base metal can tolerate without cracking:

(Eq 7) The higher the carbon equivalent of an alloy, the lower the critical or allowable cooling rate. The use of a low-hydrogen welding electrode also becomes more important. Preheating should be applied to adjust the cooling rate accordingly.

Weld Cracking Most evidence indicates that a weld-cracking failure mechanism is microstructure related. In the case of cold cracking, recent crack tip opening displacement (CTOD) results show that the reduction in toughness of HSLA weldments is related to the CGHAZ and that cracks generally propagate along or near the fusion line. The CGHAZ of microalloyed steel welds generally exhibits the highest hardness of the entire HAZ. The high-carbon untempered martensite in this region is the major cause of embrittlement. The amount of precipitates (carbides, nitrides, and carbonitrides) is found to be the highest in the regions next to the SCHAZ and the lowest at or next to the fusion line. As a result, there is a slight increase in microalloying element in solution in the CGHAZ, which increases the hardenability of this region. Hydrogen-Induced Cracking. The effect of hydrogen on weld cracking should also be mentioned. Moisture pickup from the atmosphere that is incorporated into the molten puddle, either directly or via the welding consumables, is the main source of hydrogen. The presence of hydrogen increases the HAZ cracking susceptibility of high-strength steel weldments. Also known as


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underbead, cold, or delayed cracking, it is perhaps the most serious and least understood of all weld-cracking problems. It generally occurs at the temperature below approximately 95 °C (200 °F) either immediately upon cooling or after a period of several hours. The crack can be both transgranular and intercrystalline in character, but mainly follows prior-austenite grain boundaries. The initiation of cold cracking is particularly associated with notches, such as the toe of the weld, or with inhomogeneities in microstructure that exhibit sudden changes in hardness, such as slag inclusions, martensite/ferrite interfaces, or even grain boundaries. Like most other crack growth phenomena, hydrogen-induced cracking is accentuated in the presence of high-restraint weld geometries and matrix hardening. Such cracking is associated with the combined presence of three factors: • The presence of hydrogen in the steel (even very small amounts, measured in parts per million) • A microstructure that is partly or wholly martensitic • High residual stresses (generally as a result of thick material) If one or more of these conditions is absent or at a low level, hydrogen-induced cracking will not occur. However, high cooling rates such as those found in manual processes further enhance the probability of weld HAZ cold cracking. The tolerance of steels for hydrogen decreases with increasing carbon or alloy content. Hydrogen-induced cracking can be controlled by choosing a welding process or an electrode that produces little or no hydrogen. Postweld heat treatments can be used to decrease or eliminate the residual hydrogen or to produce a microstructure that is insensitive to hydrogen cracking. Finally, welding procedures that result in low restraining stresses will also reduce the risk of weld cracking. Stress-relief cracking due to reheating is of concern when welding quenched and tempered grades and heat-resistant steels containing significant levels of carbide formers, such as chromium, molybdenum, and vanadium. When weldments of these steels are heated above approximately 510 °C (950 °F), intergranular cracking along the prior-austenite grain boundaries may take place in the CGHAZ. Also known as reheat cracking and stress-rupture cracking, stress-relief cracking is thought to be closely related to the phenomenon of creep rupture. Furthermore, during reheating, the reprecipitation of carbides is likely to occur, further increasing the hardness. The precipitation of carbides during stress relaxation alters the delicate balance between resistance to grain-boundary sliding and resistance to deformation within the coarse grains of the heat-affected zone. Some procedures that can be used singly or in combination to decrease stress-relief cracking in steels include the selection of a more appropriate weld joint design, weld location, and sequence of assembly to minimize restraint and stress concentrations. Selecting a filler metal that will provide a weld metal that has significantly lower strength than that of the HAZ at the heat-treating temperature is another way to minimize stress-relief cracking. Peening each layer of weld metal to generate a surface compressive stress state that counteracts shrinkage stresses is also very effective. Lamellar cracking, better known as lamelar tearing, is characterized by a steplike crack parallel to the rolling plane. Figure 8 shows a typical feature of lamellar tearing, the horizontal and vertical cracking of the base plate. The problem occurs particularly when making tee and corner joints in thick plates such that the fusion boundary of the weld runs parallel to the plate surface. High tensile stresses can develop perpendicular to the midplane of the steel plate, as well as parallel to it. This tearing, usually associated with inclusions in the steel, progresses from one inclusion to another. Fig. 8 Lamellar tear caused by thermal contraction strain

There is some evidence that sensitivity to lamellar tearing is increased by the presence of hydrogen in the steel. Inclusions that contain low-melting compounds, such as those of sulfur and phosphorus, also increase the sensitivity of steel to lamellar tearing by wetting the prior-austenite grain boundaries; this makes them too weak and fragile to withstand the thermal stresses during cooling. Some approaches that can minimize lamellar tearing are: • Changing the location and design of the welded joint to minimize through-thickness strains • Using a lower-strength weld metal


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• • • •


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Reducing available hydrogen Buttering the surface of the plate with weld metal prior to making the weld Using preheat and interpass temperatures of at least 95 °C (200 °F) Using base plates with inclusion shape control

Hot cracking, or solidification cracking, occurs at elevated temperatures and is usually located in the weld metal. Hot cracking also can be found in the HAZ, where it is known as liquation cracking. Solidification cracking in weld deposits during cooling occurs predominately at the weld centerline or between columnar grains. The fracture path of a hot crack is intergranular. The causes of solidification cracking are well understood. The partition and rejection of alloying elements at columnar grain boundaries and ahead of the advancing solid/liquid interface produce significant segregation. The elements of segregation form low-melting phases or eutectic structures to produce highly wetting films at grain boundaries. They weaken the structure to the extent that cracks form at the boundaries under the influence of the tensile residual stresses during cooling. Liquation cracking is also associated with grain-boundary segregation and is aggravated by the melting of these boundaries near the fusion line. These impurity-weakened boundaries tend to rupture as the weld cools because of the high residual stresses. Inclusions. Large amounts of sulfur and phosphorus are added to some steels to provide free-machining characteristics. These steels have relatively poor weldability because of hot tearing in the weld metal caused by low-melting compounds of phosphorus and sulfur at the grain boundaries. Iron oxide and iron sulfide inclusions, if present, are also harmful because of their solubility change with temperature and their propensity to precipitate at grain boundaries, contributing to low ductility, cracking, and porosity. Laminations, which are flat separations or weaknesses that sometimes occur beneath and parallel to the surface of rolled products, have a slight tendency to open up if they extend to the weld joint.

Weldability of Steels Low-carbon steels are mainly used in structural applications. Steels with less than 0.15 wt% C may harden to 30 to 40 HRC. Plain carbon steels containing less than 0.30 wt% C and 0.05 wt% S can be welded readily by most methods with little need for special measures to prevent weld cracking. The welding of sections that are more than 25 mm (1 in.) thick, particularly if the carbon content of the base metal exceeds 0.22 wt%, may require that the steel be preheated to approximately 40 °C (100 °F) and stress relieved at approximately 525 to 675 °C (1000 to 1250 °F). For low-carbon steels, a low-alloy filler metal is generally recommended for meeting mechanical property requirements. The general procedure is to match the filler with the base metal in terms of strength or, for dissimilar welds, to match the lower-strength material. Often, however, higher-strength weld metal may actually require a softer HAZ to undergo a relatively large amount of strain when the joint is subjected to deformation near room temperature. Nevertheless, a low-strength filler metal should not be used indiscriminately as a remedy for cracking difficulties. Medium-Carbon Steels. If steel containing about 0.5 wt% C is welded by a procedure commonly used for low-carbon steel, the heat-affected zone is likely to be hard, low in toughness, and susceptible to cold cracking. As indicated previously, preheating the base metal can greatly reduce the rate at which the weld area cools, thus reducing the likelihood of martensite formation. Postheating can further retard the cooling of the weld or can temper any martensite that might have formed. The appropriate preheat temperature depends on the carbon equivalent of the steel, the joint thicknesses, and the welding procedure. With a carbon equivalent in the 0.45 to 0.60 wt% range, a preheat temperature in the range of approximately 95 to 100 °C (200 to 400 °F) is generally recommended. The minimum interpass temperature should be the same as the preheat temperature. A low-hydrogen welding procedure is mandatory with these steels. Modifications in welding procedure, such as the use of a larger V-groove or of multiple passes, also decrease the cooling rate and the probability of weld cracking. Dilution can be minimized by depositing small weld beads or by using a welding procedure that provides shallow penetration. This is done to minimize carbon pickup from the base metal and the amount of hard transformation products in the fusion zone. Low heat input to limit dilution is also recommended for the first few layers in a multipass weld. High-carbon steels generally contain over 0.60 wt% C and exhibit a very high elastic limit. They are often used in applications where high wear resistance is required. These steels have high hardenability and sensitivity to cracking in both the weld metal and the HAZ. A low-hydrogen welding procedure must be used for arc welding. Preheat and postheat will not actually retard the formation of brittle high-carbon martensite in the weld. However, preheating can minimize shrinkage stresses, and postheating can temper the martensite that forms. Successful welding of high-carbon steel requires the development of a specific welding procedure for each application. The composition, thickness, and configuration of the component parts must be considered in process and consumable selections. High-strength low-alloy steels are designed to meet specific mechanical properties rather than a chemical composition. the alloy additions to HSLA steels strengthen the ferrite, promote hardenability, and help to control grain size. Weldability decreases as yield strength increases. For all practical purposes, welding these steels is the same as welding plain carbon steels that have similar carbon equivalents. Preheating may sometimes be required, but postheating is almost never required. Quenched and tempered steels are furnished in the heat-treated condition with yield strengths ranging from approximately 350 to 1000 MPa (50 to 150 ksi), depending on the composition. The base metal is kept at less than 0.22% C for good weldability. Preheating must be used with caution when welding QT steels because it reduces the cooling rate of the weld HAZ. If the cooling rate is too slow, the reaustenitized zone adjacent to the weld metal can transform either to ferrite with regions of high-carbon martensite, or to coarse bainite, of lower strength and toughness. A moderate preheat, however, can ensure against


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cracking, especially when the joint to be welded is thick and highly restrained. A postweld stress-relief heat treatment is generally not required to prevent brittle fracture in weld joints in most QT steels. Heat-treatable low-alloy steels. Examples of HTLA steels include AISI 4140, AISI 4340, AISI 5140, AISI 8640, and 300M. The high hardness of these steels requires that welding be conducted on materials in an annealed or overtempered condition, followed by heat treatment to counter martensite formation and cold cracking. However, high preheating is often used with a low-hydrogen process on these steels in a quenched and tempered condition, as in motor shaft applications. Preheating, or interpass heating, for both the weld metal and the HAZ are recommended. Hydrogen control is also essential to prevent weld cracking. Extremely clean vacuum-melted steels are preferred for welding. Low sulfur and phosphorus, as described previously, are required to reduce hot cracking. Segregation, which occurs because of the extended temperature range at which solidification takes place, reduces high-temperature strength and ductility. Fillers of lower carbon and alloy content are highly recommended. Preheat and interpass temperatures of 315 °C (600 °F) or higher are very harsh environments for welders because of the physical discomfort and because an oxide layer forms at the weld joint. However, the cooling rate must be controlled to allow the formation of a bainitic microstructure instead of the hard martensite. The bainitic microstructure can be heat treated afterward to restore the original mechanical properties of the structure. Specifications and procedures should be followed rigorously for difficult-to-weld materials. Precoated Steels. Thin plates and steel sheets are often precoated to protect them from oxidation and corrosion. The coatings commonly used are aluminum (aluminized), zinc (galvanized), and zinc-rich primers. As expected, the coating originally at the weld region is destroyed during fusion welding, and the effectiveness of the coating adjacent to the weld is significantly decreased by the welding heat. In the case of aluminized steels, the formation of aluminum oxide may adversely affect the wetting and weld pool shape. The welding electrode and filler metals should be selected carefully. A basic coating shielded metal arc (SMA) welding electrode is recommended. For galvanized steels, weld cracking is generally attributed to intergranular penetration by zinc. Zinc dissolves considerably in iron to form an intermetallic compound at temperatures close to the melting temperature of zinc. Thus, molten zinc penetrates along the grain boundaries, leaving behind a brittle champagne fracture during cooling with the onset of a tensile stress state. Cracking occurs primarily at the throat region of a fillet weld, where shrinkage strain is more significant. The use of hot-dipped coatings results in more severe cracking, while thin electrogalvanized coatings are the least susceptible to cracking. Low-silicon electrodes and rutile-base SMA welding rods are both good for galvanized steel welding. Specific welding and setup procedures should be followed, such as removing the zinc coating by an oxy-fuel process or by grinding, ensuring a large root opening, and using a slower welding speed to allow zinc vaporization and to prevent zinc entrapment in the weld metal. Adequate ventilation and fuel extraction should be mandatory in welding galvanized steels because of the health hazard of zinc fumes

Weldability Tests Weldability tests conducted to provide information on the service and performance of welds. However, the data obtained in these tests can also be applied to the design of useful structures. Frequently these data are obtained from the same type of test specimens used in determining the base metal properties. Predicting the performance of structures from a laboratory-type test is very complex because of the nature of the joint, which is far from homogeneous, metallurgically or chemically. Along with the base metal, the weld joint consists of the weld metal and the HAZ. Thus, a variety of properties are to be expected throughout the welded joint. Careful interpretation and application of the test results are required. There are currently many tests that evaluate not only the strength requirements of steel structures, but also the fracture characteristics and the effect of environmental conditions on early failure of the weldments. Selected major tests are described below. Weld Tension Test. To obtain an accurate assessment of the strength and ductility of welds, several tension test specimens can be used; all weld metal specimens, transverse weld specimens, and longitudinal weld specimens are shown in Fig. 9 . In the all weld metal test, base metal dilution should be minimized if the test is to be representative of the weld metal. However, the resulting properties may not be easy to translate into those properties achievable from welds made in an actual weld joint. Fig. 9 Typical tension test specimens for evaluating welded joints. Both plate-type specimens have identical dimensions. All dimensions given in millimeters


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Interpreting test results for the transverse butt weld test is complicated by the different strengths and ductilities generally found in the various regions of the joint. The primary information gained from the test is the ultimate tensile strength. Yield strength and elongation requirements are generally not specified. Tests of HAZ properties that are unaffected by the presence of either base metal or weld metal are not easy to conduct because it is practically impossible to obtain specimens made up entirely of the HAZ. In addition, as indicated earlier, the HAZ is composed of various regions, each with its own distinct properties. Simulated HAZ specimens that are generated and tested using a Gleeble thermomechanical testing system can be used to provide a more accurate assessment of the tensile properties of this region. Bend Test. Different types of bend tests are used to evaluate the ductility and soundness of welded joints. Bend test results are expressed in various terms, such as percent elongation in outer fibers, minimum bend radius prior to failure, go/no-go for specific test conditions, and angle of bend prior to failure. Various specimen designs, both notched and unnotched, and testing techniques have been used. Today, unnotched specimens can be used in quality control tests, while notched specimens may be used to predict in-service behavior; however, most notched bend tests are used for research purposes and are not in common industrial use. Transverse bend tests are useful because they quite often reveal the presence of defects that are not detected in tension tests. However, the transverse specimen suffers from the same weakness as the transverse weld tension test specimen in that nonuniform properties along the length of the specimen can cause nonuniform bending, although this is often compensated for by the use of a wraparound bend fixture. Hardness testing can be used to complement information gained through tension or bend tests by providing information about the metallurgical changes caused by welding. Routine methods for the hardness testing of metals are well established. In carbon and low-alloy steels, the hardness near the fusion line in the HAZ may be much higher than in the base metal because of the formation of martensite. In the HAZ areas where the temperature is low, the hardness may be lower than in the base metal because of tempering effects. The drop-weight test design is based on service failures resulting from brittle fracture initiation at a small flaw located in a region of high stress. The drop-weight test can be considered a limited-deflection bend test that uses a crack starter to introduce a running crack in the specimen. The specimen is a bar on which a brittle crack starter weld is deposited. This overlay cracks when the bar is deflected by the drop weight. A series of test is performed at different temperatures to determine the testing temperature below which the crack will propagate to the edges of the specimen. This critical temperature is also called the nilductility temperature (NDT), defined as the highest temperature at which the propagating crack reaches the edge of the specimen. Therefore, the drop-weight test is also known as the NDT test. The Charpy V-notch (CVN) test is the most popular technique for evaluating the impact properties of welds. The energy absorbed by a sample at fracture determines the toughness of the specimen. In this test, specimens at different temperatures are broken using a pendulum hammer. A typical plot of CVN results for a carbon and low-alloy steel is illustrated in Fig. 10 . The plot shows that there is a transition from low- to high-energy fracture over a narrow temperature range; this is associated with a change from trans-crystalline to ductile fracture. Therefore, material quality can be defined in terms of this transition temperature. Fig. 10 Schematic impact transition curve for steel. FATT50, fracture appearance transition temperature


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In the CVN testing of welds, the notch is typically located at the weld centerline. For CVN testing of the heat-affected zone, the notch is more typically introduced at the CGHAZ. However, because precise location of a notch is never simple in the HAZ, simulated weld samples are used instead. The crack tip opening displacement test measures toughness, primarily for elastic-plastic conditions. In CTOD tests, the clip gage opening at the onset of fracture is measured and used to calculate the crack opening displacement at the crack tip. The critical value of CTOD at fracture, ∆c, is a critical strain parameter that is analogous to the critical stress-intensity parameter. KIc. The CTOD test provides a useful method of determining the critical flaw size. Nevertheless, the test is very sensitive to changes in sample thickness, hardness, and strength, and it is difficult to obtain valid results in practical specimen thicknesses. The application of fracture mechanics to the prevention of catastrophic failure in weldments is, however, complicated by the nature of the weldment. In addition to their metallurgical heterogeneity, weldments often contain high residual stresses. Consequently, it is inadequate to fracture test the base metal and assume that the critical crack length thus determined is valid when the base metal is made into a weldment. The fracture toughness criterion must be determined for the base metal, the HAZ, and the weld metal. By first determining the zone with the lowest toughness value, it is then possible to evaluate a more realistic critical flaw size. However, the plane-strain fracture toughness tests are preproduction or pilot plant type tests that provide a rational means for designs and engineers to estimate the effects of new designs, materials, or fabrication practices on the fracture-safe performance of structures. Other popular tests include compact tension (CT) and wedge opening load (WOL) tests, which are commonly used in the evaluation of structural weldments. Further discussion of CTOD and other fracture toughness testing of welds is available in the "Selected References" at the end of this article. Stress-Corrosion Cracking Test. The presence of corrosive environments in a steel weldment may accelerate the initiation of a crack. Usually, the higher the strength of the steel, the more susceptible it becomes to stress-corrosion cracking. The steels considered in this article are not usually exposed to severely corrosive environments, but rather to the atmosphere, moisture, hydrocarbons, fertilizers, and soils. Nevertheless, welding can lower corrosive resistance by the introduction of: • Compositional differences that promote galvanic attack between weld metal, HAZ, and base metal when the joint in immersed in a conducting liquid • Residual stresses that can cause stress-corrosion cracking • Surface flaws that can act as sites for stress-corrosion cracking Stress-corrosion cracking is generally delayed cracking, with longer time to failure at lower stresses. Most stress-corrosion tests are fairly long in terms of time because of the slow crack initiation that occurs in unnotched test bars. However, it has been found that the long initiation period can be eliminated by testing precracked specimens. Additional information on the stress-corrosion cracking test is available in the "Selected References" at the end of this article.

Fabrication Weldability Tests There are various types of tests for determining the susceptibility of the weld joint to different types of cracking during fabrication. They are: • • • •

Restraint tests Externally loaded tests Underbead cracking tests Lamellar tearing tests

Table 1 summarizes the applications, controllable test variables, and typical test data of several fabrication weldability tests to illustrate the differences among them. Of the many tests identified in Table 1 , the Lehigh restraint test, the Varestraint test, and the controlled thermal severity test are described below. Table 1 Comparison of weldability tests for fabrication


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Test


Fields of use

Controllable variables

01 Sep 2005

Type of data

Specific equipment

Cost

Lehigh restraint test

Weld metal hot and cold cracks, root cracks, HAZ hydrogen cracks, stress-relief cracks

Critical restraint, or % Joint geometry, process, hindered control filler metal, restraint level, heat input, preheat, postweld heat treatment

None

Costly machining

Slot test

HAZ hydrogen cracks

Filler, metal, interpass time, Time to crack, critical preheat preheat

None

Low cost

Rigid restraint test

Weld metal hot and cold cracks, root cracks, HAZ hydrogen cracks

Joint geometry, process restraint level, filler metal, heat input, preheat

Restraint jig

Costly machining and setup

Tekken test

Weld metal root cracks, HAZ hydrogen cracks

Joint geometry, process filler Critical preheat metal, heat input, preheat

None

Low cost

Circular groove test

Weld metal hot and cold cracks, HAZ hydrogen cracks

Process, filler metal, preheat Go/no-go

None

Costly preparation

Implant test

HAZ hydrogen cracks, stress-relief cracks

Process, filler metal, preheat, Critical fracture stress, Loading jig postweld heat treatment critical preheat

Intermediate cost

Tension restraint cracking test

HAZ hydrogen cracks

Process, filler metal, heat input, preheat

Critical fracture stress, Loading jig critical preheat

Costly machining and setup

Varestraint test

Weld metal and HAZ hot cracks

Process, filler metal, heat input

Crack length, % strain

Loading jig

Costly preparation and analysis

Longitudinal bead-on-plate test

HAZ hydrogen cracks

Electrical type, heat input

% cracking

None

Low cost

Controlled thermal severity test

HAZ hydrogen cracks in fillet welds

Electrical type, cooling rate, Go/no-go (at two preheat cooling rates)

None

Costly preparation

Cruciform test

HAZ hydrogen cracks, weld Process, heat input, preheat, Go/no-go metal root cracks filler metal

None

Costly preparation

Lehigh cantilever test

Lamellar tearing

Process, filler metal, heat input, preheat

Critical restraint stress Loading jig and strain

Costly specimen preparation

Cranfield test

Lamellar tearing

Filler metal

Number of passes to crack

None

Low cost

Weld metal soundness

Filler metal

Go/no-go

None

Low cost

Nick bend test Source: Ref 1

Critical restraint

The Lehigh restraint test (Fig. 11 ) is particularly useful for quantitatively rating the crack susceptibility of a weld metal as affected by electrode variables. This test provides a means of imposing a controllable severity of restraint on the root bead that is deposited in a butt weld groove with dimensions suitable to the application. Slots are cut in the sides and ends of a plate prior to welding. By changing the length of the slots, the degree of plate restraint on the weld is varied without significantly changing the cooling rate of the weld. Therefore, a critical restraint for cracking can be determined for given welding conditions. This sample is also useful for hydrogen cracking. Fig. 11 Basic outline of a Lehigh test specimen


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The Varestraint test (Fig. 12 ) determines the susceptibility of the welded joint to hot cracking. The test utilizes external loading to impose controlled plastic deformation in a plate while a weld bead is being deposited on the long axis of the plate. The specimen is mounted as a cantilever beam, and a pneumatically driven yoke is positioned to force the specimen downward when the welding arc reaches a predetermined position. By the choice of the radius to which the plate is bent, the severity of deformation causing cracking can be determined. Strain from 0 to 4% can be chosen according to the susceptibility of the joint to hot cracking. When the bending moment is applied transverse to the weld axis, the test is termed transvarestraint. A spot Varestraint test can also be conducted by keeping the arc stationary; bending is applied at the moment the arc is extinguished. Fig. 12 Schematic of the Varestraint test. A, weld location; B, die; C, arc; D, load; r, radius of deformation

The controlled thermal severity test (Fig. 13 ) is designed to measure the cracking sensitivity of steels under cooling rates controlled by the thickness of the plates and the number of paths available for dissipating the welding heat. It is conducted with a plate bolted and anchor welded to a second plate in a position to provide two fillet (lap) welds. The fillet located at the plate edges has two paths of heat flow. The lap weld located near the middle of the bottom plate has three paths of heat flow, thus inducing faster cooling. The fillet welds are made first and allowed to cool, followed by the lap welds. After a holding time of 72 h at room temperature, the degree of cracking is determined by measuring the crack length on metallographic specimens. Fig. 13 Schematic of the controlled thermal severity test


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A number of other tests have been developed that contain welds in a circular configuration. The circular patch test has probably the most severe testing conditions; the two varieties are the Navy circular patch restraint test and the segmented circular patch restraint level. Cracking is detected by visual, radiographic, and liquid penetrant inspection. The cracking susceptibility of a material is measured as the total crack length and expressed as a percentage of the weld length. These tests can be used to determine both hot and cold cracking in the weld metal and the HAZ. Depending on the results, a go/no-go criteria is established for weld qualification. Detailed descriptions of these tests can be found in the "Selected References" at the end of this article. REFERENCE 1. R. Stout, Weldability of Steels, 4th ed., Welding Research Council, 1987 SELECTED REFERENCES • R.R. Barr and F.M. Burdekin, Design Against Brittle Failure, in Rosenhain Centenary Conference Proceedings, R.G. Baker and A. Kelly, Ed., The Royal Society, p 85, 1975 • O. Blodgett, "Why Preheat? An Approach to Estimating Correct Preheat Temperature," Brochure G-231, Lincoln Arc Welding Foundation, June 1970 • B.F. Brown, "Stress Corrosion Cracking and Corrosion Fatigue of High Strength Steels," Report 210, Defense Metals Information Center, 1964, p 91−102 • "Classification of Microstructures in Low Carbon Low Alloy Steel Weld Metal and Terminology," DOC IX-1282-83, International Institute of Welding, 1983 • J. Cornu, Advanced Welding Systems: Fundamentals of Fusion Welding Technology, IFS/Springer Verlag, 1988 • C.L.M. Cottrell, Hardness Equivalent May Lead to a More Critical Measure of Weldability, Met. Constr., Vol 16 (No. 12), 1984, p 740−743 • C.E. Cross, Ø. Grong, S. Liu, and J.F. Capes, Metallography and Welding Process Control, in Applied Metallography, G. Vander Voort, Ed., Van Nostrand Reinhold, 1985 • G.J. Davies and J.G. Garland, Solidification Structures and Properties of Fusion Welds, Int. Met. Rev., No. 20, 1975, p 83−106 • K. Easterling, Introduction to the Physical Metallurgy of Welding, Butterworths, 1983


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• D.P. Fairchild, Brittle Zones in Structural Welds, in Welding Metallurgy of Structural Steels, J. Koo, Ed., The Metallurgical Society, 1987, p 303−318 • H. Granjon, "Notes on the Carbon Equivalent," DOC IX-555-67, International Institute of Welding, 1967 • J.D. Harrison, M.G. Davies, G.L. Archer, and M.S. Kamath, "The COD Approach and Its Application to Welded Structures," Report 55/1978, The Welding Institute, 1978 • Y. Ito and K. Bessyo, "Weldability Formula of High Strength Steels Related to Heat Affected Zone Cracking," DOC IX-576-68, International Institute of Welding, 1968 • J. Koo and A. Ozekan, Local Brittle Zone Microstructure, in Welding Metallurgy of Structural Steels, J. Koo, Ed., The Metallurgical Society, 1987, p 119−135 • S. Kou, Welding Metallurgy, Wiley Interscience, 1987 • S. Liu, D.L. Olson, and D.K. Matlock, A Thermodynamic and Kinetic Approach in the Development of Expressions for Alloy Behavior Prediction, J. Heat Treat., Vol 4 (No. 4), 1986, p 309−316 • F. Matsuda, T. Hashimoto, and T. Senda, Fundamental Investigations on Solidification Structure, Trans. Natl. Res. Inst. Met. (Jpn.), Vol 11, 1969, p 43−58 • H.G. Pisarski and J. Kudoh, Exploratory Studies on the Fracture Toughness of Multi-Pass Welds With Locally Embrittled Regions, in Welding Metallurgy of Structural Steels, J. Koo, Ed., The Metallurgical Society, 1987, p 263−276 • Properties and Selection: Irons and Steels, Vol 1, 9th ed., Metals Handbook, American Society for Metals, 1978 • S.T. Rolfe, "Development of a KIc Stress Corrosion Specimen," Technical Report, United States Steel Applied Research Laboratory, 1965 • A.B. Rothwell, CAN/MET Report 79-6, Can. Weld. Fabr., Vol 20, 1980 • C.P. Royer, A User's Perspective on HAZ Toughness, in Welding Metallurgy of Structural Steels, J. Koo, Ed., The Metallurgical Society, 1987, p 255−262 • H. Suzuki "Carbon Equivalent and Maximum Hardness," DOC IX-1279-83, International Institute of Welding, 1983 • Welding, Brazing, and Soldering, Vol 6, 9th ed., Metals Handbook, American Society for Metals, 1983 • Welding Handbook, Vol I and II, 7th ed., American Welding Society, 1983


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Elevated-Temperature Properties of Ferri...

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Service Characteristics of Carbon and Low-Alloy Steels Elevated-Temperature Properties of Ferritic Steels CARBON STEELS and low-alloy steels with ferrite-pearlite or ferrite-bainite microstructures are used extensively at elevated temperatures in fossil-fired power-generating plants, aircraft power plants, chemical-processing plants, and petroleum-processing plants. Carbon steels are often used up to about 370 °C (700 °F) under continuous loading, but also have allowable stresses defined up to 540 °C (1000 °F) in Section VIII of the ASME Boiler and Pressure Vessel Code. Carbon-molybdenum steels with 0.5% Mo are used up to 540 °C (1000 °F), while low-alloy with 0.5−1.0% Mo in combination with 0.5−9.0% Cr and sometimes other carbide formers (such as vanadium, tungsten, niobium, and titanium) are often used up to about 650 °C (1200 °F). For temperatures above 650 °C (1200 °F), austenitic alloys are generally used. However, these general maximum-use temperature limits do not necessarily apply in specific applications with different design criteria. Tables 1 and 2 , for example, list maximum-use temperatures in two specific application areas with different design criteria. Table 1 Temperature limits of superheater tube materials covered in ASME Boiler Codes Maximum-use temperature Oxidation/graphitization criteria, metal surface(a)

Strength criteria, metal midsection

°C

°F

°C

°F

400−500

750−930

425

795

Ã0.5Cr-0.5Mo

550

1020

510

950

Ã1.2Cr-0.5Mo

565

1050

560

1040

Ã2.25Cr-1Mo

580

1075

595

1105

Ã9Cr-1Mo

650

1200

650

1200

Material SA-106 carbon steel Ferritic alloy steels

Austenitic stainless steel ÃType 304H 760 1400 815 1500 (a) In the fired section, tube surface temperatures are typically 20−30 °C (35−55 °F) higher than the tube midwall temperature. In a typical U.S. utility boiler, the maximum metal surface temperature is approximately 625 °C (1155 °F).

Table 2 Suggested maximum temperatures in petrochemical operations for continuous service based on creep or ruptured data Maximum temperature based on rupture

Maximum temperature based on creep rate Material

°C

°F

°C

°F

Carbon steel

450

850

540

1000

C-0.5 Mo steel

510

950

595

1100

21=4 Cr-1Mo steel

540

1000

650

1200

Type 304 stainless steel

595

1100

815

1500

Alloy C-276 nickel-base alloy

650

1200

1040

1900

This article covers some elevated-temperature properties of carbon steels and low-alloy steels with ferrite-pearlite and ferrite-bainite microstructures for use in boiler tubes, pressure vessels, and steam turbines. In these applications, the selection of steels to be used at elevated temperatures generally involves compromise between the higher efficiencies obtained at higher


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operating temperatures and the cost of equipment, including materials, fabrication, replacement, and downtime costs. The highly alloyed steels, which depend on an austenitic matrix for their high-temperature properties, generally have higher resistance to mechanical and chemical degradation at elevated temperatures than the low-alloy ferritic steels. However, a higher alloy content generally means higher cost. Therefore, carbon and low-alloy ferritic steels are extensively used in several forms (piping, pressure vessel plates, bolts, structural parts) in a variety of applications that involve exposure to elevated temperatures. In addition, interest in ferritic steels has increased recently because their relatively lower thermal expansion coefficient and higher thermal conductivity make them more attractive than austenitic steels in applications where thermal cycling is present. To illustrate the tonnage requirements for carbon and low-alloy steels in industrial construction, 1360 Mg (1500 tons) of pressure tubing were required for the construction of a single 500 MW coal-fired generating plant. The quantities of the various carbon and low-alloy steels used in the pressure tubing were as follows: Steel type

Tons

% of total tonnage

540

36

Carbon 1

C- =2Mo

150

10

11=4Cr-1Mo

495

33

21=4Cr-1Mo

150

10

9Cr-1Mo

165

11

This list of carbon and low-alloy steels is for pressure tube applications and does not include the chromium-molybdenum-vanadium steels that are used for turbine rotors, high-temperature bolts, and pressure tubing.

Carbon and Low-Alloy Steels for Elevated-Temperature Service The numerous types of steels used in elevated-temperature applications include the following: • • • • • •

Carbon steels Low-alloy steels High steels Stainless steels Hot-work tool steels Iron-base superalloys

Within the context of this article, the low-alloy steels considered are the creep-resistant steels with 0.5 to 1.0% Mo combined with 0.5 to 9.0% Cr and perhaps other carbide formers (such as vanadium, tungsten, niobium, and titanium). High-strength low-alloy (HSLA) steels are not considered here because they typically have molybdenum contents below 0.5%, which limits their resistance against creep and temper embrittlement. However, HSLA steels, which are discussed in the article "High-Strength Structural and High-Strength Low-Alloy Steels" in this Volume, may be effective substitutes for carbon steels in elevated-temperature applications. Another category of ferritic steels for elevated-temperature service are manganese-molybdenum-nickel ferritic steels (ASTM A 302 and A 533), which are commonly used for pressure vessels in light-water reactors. High-alloy steels, stainless steels, hot-work tool steels, and the iron-base superalloys are discussed in the Section "Specialty Steels and Heat-Resistant Alloys" in this Volume. Alloy Designations and Specifications Carbon and low-alloy steels used for elevated-temperature service are usually identified by American Iron and Steel Institute (AISI) designations; aerospace material specification (AMS), American Society of Mechanical Engineers (ASME), or American Society for Testing and Materials (ASTM) specification number; nominal composition; or trade name. These steels have also been assigned numbers in the Unified Numbering System. In addition, there are Military and Federal specifications covering many of these steels. Steel products manufactured for use under the ASME Boiler and Pressure Vessel Code must comply with provisions of the appropriate ASME specification. Each specification includes information on ranges and limits of composition, dimensions and tolerances, minimum mechanical properties, and other functional requirements. The designations applied to these products include the letters "SA," the number of the specification, and possibly other letters or numbers to distinguish among the various types, grades, and classes within a single specification. Most ASTM specifications are identical to the ASME specification of the same number except that the ASTM designations begin with the letter "A." Some examples of ASME specifications for elevated-temperature steels, as well as their compositions and typical room-temperature mechanical properties, are given in Tables 3(a) and 3(b) . Table 3(a) Compositions of steels for elevated-temperature service ASME

UNS

Nominal


Composition, %

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specificatio designa composit n tion ion


Product form

C

Mn

Si

P

S

01 Sep 2005

Cr

Ni

Mo

Others

SA-106A

K02501

C

Seamless carbon steel pipe

0.25(a) 0.27−0. 0.10(b) 0.048 0.058(a (a) ) 93

...

...

...

...

SA-106B

K01700

C-Si

Seamless carbon steel pipe

0.30(a) 0.29−1. 0.10(b) 0.048 0.058(a (a) ) 06

...

...

...

...

SA-285A

K03006

C

Carbon steel PV plate

0.17(a) 0.90(a)

0.035 0.045(a (a) )

...

...

...

0.25 Cu(a)

SA-299

K02803 C-Mn-Si C-Mn-Si steel PV plate

0.28(a) 0.90−1. 0.15−0. 0.035 0.040(a (a) ) 40 30

...

...

...

...

SA-204A

K11820 C-1=2Mo Mo alloy steel PV plate

0.18(a) 0.90(a) 0.15−0. 0.035 0.040(a (a) ) 30

...

...

0.45−0. 60

...

SA-302A

K12021 Mn-Mo Mn-Mo-Mn and Mo-Ni alloy PV plate

0.20(a) 0.95−1. 0.15−0. 0.035 0.040(a (a) ) 30 30

...

...

0.45−0. 60

...

0.25(a) 1.15−1. 0.15−0. 0.035 0.040(a (a) ) 50 30

...

SA-533B2 K12539 Mn-Mo- Mn-Mo-Mn and Ni Mo-Ni alloy steel PV plate ...

...

SA-517F

K11576

SA-335 P12

0.15(a) 0.30−0. 0.50(a) 0.045 0.045(a 0.50−1. K11562 1Cr-1=2M Seamless ferritic (a) ) alloy steel pipe for o 61 25 high-temperature service

0.40−0. 0.45−0. 70 60

0.10 Cu(a)

High-strength alloy 0.10−0. 0.60−1. 0.15−0. 0.035 0.040(a 0.40−0. 0.70−1. 0.40−0. 0.002−0.006 steel PV plate (a) ) 20 00 35 65 00 60 B, 0.15−0.050 Cu, 0.03−0.08 V ...

0.44−0. 65

...

SA-217WC J12072 11=4Cr-1=2 Alloy steel 6 castings Mo

0.20(a) 0.50−0. 0.60(a) 0.04( 0.045(a 1.00−1. a) ) 80 50

...

0.45−0. 65

...

SA-387Gr2 K21590 21=4Cr-1 Cr-Mo alloy steel 2 PV plate Mo

0.15(a) 0.30−0. 0.50(a) 0.035 0.035(a 2.0−2.5 (a) ) 60

...

0.90−1. 10

...

SA-387Gr5 S50100 5Cr-1=2M Cr-Mo alloy steel PV plate o

0.15(a) 0.30−0. 0.50(a) 0.040 0.030(a 4.0−6.0 (a) ) 60

...

0.45−0. 65

...

0.02(a) 0.35−0. 1.00(a) 0.04( 0.045(a 8.0−1.0 a) ) 65

...

0.90−1. 20

...

SA-217C12 J82090 9Cr-1Mo Alloy steel castings (a) Maximum. (b) Minimum

Table 3(b) Room-temperature mechanical properties of steels for elevated-temperature service listed in Table 3(a) Mechanical properties Tensile strength

Yield strength, minimum

Minimum elongation in 50 mm (2 in.), %

Minimum reduction in area, %

ASME specification

MPa

ksi

MPa

ksi

SA-106A

330

48(a)

207

30

35(b), 25(c)

...

SA-106B

415

60(a)

241

35

30(b), 16.5(c)

...

SA-285A

310−380

45−55

165

24

27(d), 30

...

SA-299

515−620

75−90

290

42

16(d)

...

37

19(d), 23

...

SA-204A

445−530

65−77

255

SA-302A

515−655

75−95

310

45

15(d), 19

...

SA-533B2

620−790

90−115

475

70

16

...

SA-517F

795−930

115−135

690

100

16

34−45

SA-335P12

415

60(a)

207

30

30(b), 20(c)

...

SA-217WC6

485−620

70−90

275

40

20

35

SA-387Gr22-1

415−585

60−85

207

30

18(d), 45

40

SA-387Gr5-2

515−690

75−100

310

45

18(d), 22

45

90−115

415

60

18

35

SA-217C12

620−795


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01 Sep 2005

(a) Minimum. (b) Longitudinal. (c) Transverse. (d) Elongation in 200 mm (8 in.)

Aerospace material specifications, as the name suggests, are specifications for products intended for the aerospace industry. The nominal compositions, typical applications, and typical mechanical properties of steels often identified by AMS numbers are given in Table 4 . Table 4 Compositions and mechanical properties of AISI steels for elevated-temperature service AISI designatio n

Nominal composition, % AMS Commercial UNS designations designation designations

Typical applications

C

Mn

Si

Cr

Mo

V

...

K14675

Bolting and structural parts

0.46

0.60

0.26

1.00

0.50

0.30

6302, 6385, 6458

17−22 AS

K23015

Bolting and structural parts

0.30

0.55

0.65

1.25

0.50

0.25

603

6303, 6436

17−22 AV

K22770

Turbine rotors and aircraft parts

0.27

0.75

0.65

1.25

0.50

0.85

610

6437, 6485

H11 mod

T20811 K74015

Ultrahigh-strength components

0.40

0.30

0.90

5.00

1.30

0.50

601

6304

602

Temperature at which 70 MPa (10 ksi) will cause rupture in

Room-temperature tensile properties AISI designatio n 601 602

Yield strength

Tensile strength

MPa

ksi

MPa

ksi

710

103

855

124

Elongation in 50 Reduction mm (2 in.), % in area, %

745−93 108−135 880−106 128−154 0 0

1000 h °C

°F

Temperature to produce min creep rate at 70 MPa (10 ksi)

10,000 h

1 µm/m · h

°C

°C

°F

°F

0.1 µm/m · h °C

°F

29

61

620

1150

595

1100

...

...

...

...

16−21

53−63

625

1160

590

1090

...

...

555

1030

603

1000

145

1100

160

17

52

650

1200

613

1135

...

...

565

1050

610

1480

215

1805

262

10

36

630

1170

595

1100

560

1040

540

1000

The AISI designation for steels intended for elevated-temperature service is a three-digit number beginning with a 6, such as 601. The AISI designations are also included in Table 4 . Carbon Steels Carbon steels are the predominant materials in pressure vessel fabrications because of their low cost, versatile mechanical properties, and availability in fabricated forms. They are the most common materials used in noncorrosive environments in the temperature range of −29 to 425 °C (−20 to 800 °F) in oil refineries and chemical plants. Although the ASME code gives allowable stresses for temperatures greater than 425 °C (800 °F), it also notes that prolonged exposure at these temperatures may result in the carbide phase of the carbon steel being converted to graphite. This phenomenon, known as graphitization, is a cumulative process dependent on the time the material is at or above 425 °C (800 °F). The result is a weakening of the steel after high-temperature exposure (Fig. 1 ). Carbon steels are also increasingly affected by creep at temperatures above 370 °C (700 °F). Figure 2 shows the effect of temperature on the stress-to-rupture life of a carbon steel. Fig. 1 Effect of elevated-temperature exposure on the room-temperature tensile properties of normalized 0.17% C steel after exposure (without stress) to indicated temperature for 83,000 h


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Fig. 2 Effect of exposure to elevated temperature on stress-to-rupture of carbon steel. Stress-to-rupture in 1000 and 10,000 h at the indicated temperature for specimens of normalized 0.17% C steel exposed to the test temperature (without stress) for 83,000 h and for similar specimens not exposed to elevated temperature prior to testing

Creep-Resistant Low-Alloy Steels Creep-resistant low-alloy steels usually contain 0.5 to 1.0% Mo for enhanced creep strength, along which chromium contents between 0.5 and 9% for improved corrosion resistance, rupture ductility, and resistance against graphitization. Small additions of carbide formers such as vanadium, niobium, and titanium may also be added for precipitation strengthening and/or grain refinement. The effects of alloy elements on transformation hardening and weldability are, of course, additional factors. The three general types of creep-resistant low-alloy steels are chromium-molybdenum steels, chromium-molybdenum-vanadium steels, and modified chromium-molybdenum steels. Chromium-molybdenum steels are used primarily for tube, pipe, and pressure vessels, where the allowable stresses may permit creep deformation up to about 5% over the life of the component. Typical creep strengths of various chromium-molybdenum steels are shown in Fig. 3 . Figure 3 also shows the creep strength of a chromium-molybdenum steel with vanadium additions. Chromium-molybdenum-vanadium steels provide higher creep strengths and are used for high-temperature bolts, compressor wheels, or steam turbine rotors, where allowable


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01 Sep 2005

stresses may require deformations less than 1% over the life of the component. Fig. 3 General comparison of creep strengths of various creep-resistant low-alloy steels

Chromium-molybdenum steels are widely used in oil refineries, chemical industries, and electrical power generating stations for piping, heat exchangers, superheater tubes, and pressure vessels. The main advantage of these steels is the improved creep strength from molybdenum and chromium additions and the enhanced corrosion resistance from chromium. The creep strength of chromium-molybdenum steels is derived mainly from two sources: solid-solution strengthening of the matrix ferrite by carbon, molybdenum, and chromium: and precipitation hardening by carbides. Creep strength generally, but not always, increases with higher amounts of molybdenum and chromium. The effects of chromium and molybdenum on creep strength are complex (see "Effects of Composition" in this article). In Fig. 3 , for example, 2.25Cr-1Mo steel has a higher creep strength than 5Cr-0.5Mo steel. Chromium-molybdenum steels are available in several product forms (see Table 24 in the article "Classification and Designation of Carbon and Low-Alloy Steels" in this Volume). In actual applications, boiler tubes are used mostly in the annealed condition, whereas piping is used mostly in the normalized and tempered condition. Bend sections used in piping, however, are closer to an annealed condition than to a normalized condition. As a result of the cooling rates employed in these treatments, the microstructures of chromium-molybdenum steels may vary from ferrite-pearlite aggregates to ferrite-bainite aggregates. Bainite microstructures have better creep resistance under high-stress, short-time conditions but degrade more rapidly at high temperatures than pearlitic structures. As a result, ferrite-pearlite material has better intermediate-term, low-stress creep resistance. Because both microstructures will eventually spheroidize, it is expected that over long service lives the two microstructures will converge to similar creep strengths. The 0.5Mo steel with 0.15% C is used for piping and superheater tubes operating at metal temperatures to 455 °C (850 °F). Above this temperature, spheroidization and graphitization may increase the possibility of failure in service. Use of carbon-molybdenum steel has been largely discontinued for the higher temperatures because of graphitization. Chromium steels are highly resistant to graphitization and are therefore preferred for service above 455 °C (850 °F). The 1.0Cr-0.5Mo steel is used for piping, cracking-still tubes, and boiler tubes for service temperatures to 510 or 540 °C (950 or 1000 °F). The similar 1.25Cr-0.5Mo steel is used up to 590 °C (1100 °F) and has comparable stress-rupture and creep properties as that of the 1.0Cr-0.5Mo alloy (Fig. 4 ). Fig. 4 Creep strength (0.01% 1000 h) and rupture strength (100,000 h) of 1Cr-0.5Mo and 1.25Cr-0.5Mo steel. Source: Ref 1


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01 Sep 2005

The 2.25Cr-1.0Mo steel has better oxidation resistance and creep strength than the steels mentioned above. The 2.25Cr-1Mo steel is a highly favored alloy for service up to 650 °C (1200 °F) without the presence of hydrogen or 480 °C (900 °F) in a hydrogen environment. This steel, which has substantial documentation of its elevated-temperature properties (Ref 2, 3, 4, 5, and 6), is discussed in more detail in the section "Elevated-Temperature Behavior of 2.25Cr-1Mo Steel" in this article. The 5, 7, and 9% Cr steels are generally lower in stress rupture and creep strength that the lower-chromium steels because the strength at elevated temperatures typically drops off with an increase in chromium. However, this may not always be the case, depending on the service temperature (Fig. 5 ) and the exposure (Fig. 6 and 7 ). Heat treatment is also an important factor. The main advantage of these steels is the improved oxidation resistance from the increased chromium content. Fig. 5 Variation of 105-h creep-rupture strength as a function of temperature for 21=4Cr-1Mo steel, standard 9Cr-1Mo, modified 9Cr-1Mo, and 304 stainless steel. Source: Ref 7


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01 Sep 2005

Fig. 6 Effect of elevated-temperature exposure on stress-rupture behavior of (a) normalized and tempered 21=4Cr-1Mo steel and (b) annealed 9Cr-1Mo steel. Exposure prior to stress-rupture testing was at the indicated test temperatures (without stress) and was 10,000 h long for the 21=4Cr-1Mo steel and 100,000 h long for the 9Cr-1Mo steel. n/a, data not available at indicated exposure and rupture life.


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01 Sep 2005

Fig. 7 Effect of temperature exposure on the room-temperature properties of (a) normalized (900 °C, or 1650 °F) and tempered (705 °C, or 1300 °F) 21=4Cr-1Mo steel after exposure (without stress) to indicated temperature for 10,000 h and (b) annealed 9Cr-1Mo steel after exposure (without stress) to indicated temperatures for 100,000 h

The chromium-molybdenum-vanadium steels are manufactured with higher carbon ranges (such as 0.28 to 0.33% and 0.40 to 0.50%) and are used in the normalized and tempered or quenched and tempered condition. Because of the relatively high yield strengths (Fig. 8 ) and creep strengths (Fig. 3 ), these steels are suitable for bolts, compressor wheels in gas turbines and steam turbine rotors, and other parts operating at temperatures up to 540 °C (1000 °F). The most common low-alloy composition contains 1% Cr, 1% Mo, and 0.25% V. Fig. 8 Room-temperature and short-time elevated-temperature tensile strengths and yield strengths of selected steels containing less than 10% alloy. The 1.0Cr-0.5Mo steel, 0.5Mo steel, type 502, and 2.25Cr-1.0Mo steel were annealed at 843 °C (1550 °F). The 1.25Cr-0.5Mo steel was annealed at 815 °C (1500 °F). The 7.0Cr-0.5Mo and 9.0Cr-1.0Mo steels were annealed at 900 °C (1650 °F). The 1.0Cr-1.0Mo-0.25V steel was normalized at 955 °C (1750 °F) and tempered at 650 °C (1200 °F). H11, hardened 1010 °C (1850 °F), tempered 565 °C (1050 °F)


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Bolt Applications. The basic compositions of low-alloy high-temperature bolt steels have evolved from chromium-molybdenum steels to chromium-molybdenum-vanadium steels. The chromium-molybdenum steels used until the late 1940s had creep strengths adequate for service at temperatures up to about 480 °C (895 °F). With the increasing need for a higher-strength steel, a 1Cr-1Mo-1=4V steel strengthened by stable V4C3 precipitates was developed. This alloy was found to be adequate for steam temperatures up to 540 °C (1000 °F). When the steam temperatures reached about 565 °C (1050 °F) in the mid-1950s, a 1Cr-1Mo-3=4V steel, in which vanadium and carbon had been stoichiometrically optimized to get the largest volume fraction of V4C3 and hence the highest creep strength, was developed. Unfortunately, this development had overlooked the importance of rupture ductility, and many creep-rupture failures of bolts due to notch sensitivity occurred. The loss in rupture ductility was subsequently countered by grain refinement and by compositional modifications involving titanium and boron. Melting practice is also another factor in improving rupture ductility. High-Temperature Rotor Applications. Since it was introduced in the 1950s, 1Cr-1Mo-0.025V steel has remained the industry standard in turbine rotor applications, although a few higher-alloy rotor steels (12% Cr) have been developed (see the article "Elevated-Temperature Properties of Stainless Steels" in this Volume). It is well recognized that 1Cr-1Mo-0.25V rotor steels are limited by their creep strength for service up to about 540 °C (1000 °F). The desired properties in chromium-molybdenum-vanadium steel rotors is made possible careful control of heat treatment and composition. In the United States, the usual practice has been to air cool the rotors from the austenitizing temperature in order to achieve a highly creep-resistant, but somewhat less tough, upper bainitic microstructure. In Europe, however, manufacturers have resorted to oil quenching of rotors from the austenitizing temperature to achieve a better compromise between creep strength and toughness. Oil quenching of 1Cr-1Mo-0.25V rotors may shift the transformation product increasingly toward lower bainite, but it is unlikely that the cooling rates needed for formation of martensite (that is, 10,000 °C/h or 20,000 °F/h) are ever encountered. Comparative evaluation of creep properties of chromium-molybdenum-vanadium steels with martensite, bainite, and ferrite-pearlite as the principal microstructure have been conducted by numerous investigators, and the results have been reviewed elsewhere (Ref 8). There is consensus that upper bainitic structures provide the best creep resistance coupled with adequate ductility. Toughness properties are discussed in Ref 9. Turbine Casing Applications. Chromium-molybdenum-vanadium steels are also used for turbine casings. The table below compares the maximum application temperatures of various low-alloy steels used for turbine casings (Ref 10): Maximum application temperature Casing material

°C

°F

C-1=2 Mo (0.25C max, 0.20−0.50Si, 0.5−1.0Mn, 0.50−0.70Mo)

480

895

Cr-1=2Mo (0.15C max, 0.60Si max, 0.5−0.8Mn, 1.0−1.5Cr, 0.45−0.65Mo)

525

975

2 =4Cr-1Mo (0.15C max, 0.45Si max, 0.4−0.8Mn)

540

1000

Cr−Mo−V (0.15C max, 0.15−0.30Si max, 0.4−0.6Mn, 0.7−1.2Cr, 0.7−1.2Mo, 0.25−0.35V)

565

1050

1

565

1050

1

=2Cr-Mo-V (0.1−0.15C, 0.45Si max, 0.4−0.7Mn, 0.4−0.6Cr, 0.4−0.6Mo, 0.22−0.28V)

Modified Chromium-Molybdenum Steels. To achieve higher process efficiencies in future coal conversion plants, chemical-processing plants, and petrochemical-refining plants, several modified versions of chromium-molybdenum pressure vessel steels have been investigated for operation at higher temperatures and pressures than those currently encountered. The higher temperatures affect the elevated-temperature strength, the dimensional deformation, and the metallurgical stability of an alloy, while higher operating pressures require either higher-strength alloys or thicker sections. Of the unmodified ferritic steels, SA-387 grade 22, class 2 (normalized and tempered 21=4Cr-1Mo unmodified steel) meets the


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requirements for the fabrication of large pressure vessels per Section VIII, Division 2 of the ASME Boiler and Pressure Vessel Code. Unfortunately, the thick-section hardenability is insufficient to prevent the formation of cementite, even with accelerated cooling procedures and lower tempering conditions (Ref 11). This persistence of cementite in thick-section SA-387 grade 22, class 2 is a concern regarding hydrogen attack (Ref 12). Other unmodified chromium-molybdenum steels such as 3Cr-1Mo and 5Cr-0.5Mo (SA-387 grades 21 and 5) resist hydrogen attack, but the design allowables are below those of 2 1=4Cr-1Mo steel at some temperatures of interest. Higher-chromium alloys, such as 7Cr-0.5Mo and 9Cr-1Mo, also have strengths below normalized and tempered 21=4Cr-1Mo and so have not been considered in the United States for heavy-wall vessels. Therefore, several modified chromium-molybdenum alloys have been investigated for thick-section vessels in a hydrogen environment. These modified chromium-molybdenum alloys contain various microalloying elements such as vanadium, niobium, titanium, and boron. Three categories (Ref 11) of modified chromium-molybdenum steels investigated for thick-section applications in a hydrogen environment are: • 3Cr-1Mo modified with vanadium, titanium, and boron (Ref 13): This steel is approved for service up to 455 °C (850 °F), is fully hardenable, resists hydrogen attack, and has strengths capable of meeting the design allowables of normalized and tempered 21=4Cr-1Mo steel • 9Cr-1Mo steel modified with vanadium and niobium (Ref 14, 15): This steel has strengths exceeding those of 2.25Cr-1Mo and is approved for use to more than 600 °C (1110 °F) for steam and hydrogen service • 21=4Cr-1Mo steel modified with vanadium, titanium, and boron (Ref 4, 16): Vanadium-modified 21=4Cr-1Mo steel is fully hardenable, resists hydrogen attack, and exceeds the strength of normalized and tempered 21=4Cr-1Mo steel Other modified alloys, such as 3Cr-1.5Mo-0.1V-0.1C, have also been investigated (Ref 13, 17). The modified alloys have improved hardenability over unmodified 21=4Cr-1Mo steel. However, these modified chromium-molybdenum steels with bainitic microstructures undergo a strain softening (Ref 18, 19, 20, 21), which may be a limitation in applications with cyclic stresses. The modified 9Cr-1Mo steel is an attractive alloy because it has strengths (Fig. 5 ) capable of meeting or exceeding the allowable stresses of stainless steel (Fig. 9 ). Microstructural work has indicated that the improved strength of the modified alloy derives from two factors. First, fine M23C6 precipitate particles nucleate on Nb(C,N), which first appears during the heat treatment. Second, the vanadium enters M23C6 and retards its growth at the service temperature. The finer distribution of M23C6 adds to the strength, and its retarded grain-size growth holds the strength for long periods of time at the service temperature. The grain-coarsening behavior of the modified 9Cr-1Mo steel as a function of normalizing temperature and time-temperature exposure is shown in Fig. 10 . Fig. 9 Estimated design allowable stresses (Section VIII of ASME Boiler and Pressure Vessel Code) as a function of temperature for modified 9Cr-1Mo steel, standard 9Cr-1Mo, 21=4Cr-1Mo steel, and 304 stainless steel. Source: Ref 7


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Fig. 10 Grain-coarsening behavior of a modified 9Cr-1Mo steel (9Cr-1Mo steel with 0.06 to 0.10% Nb and 0.18 to 0.25% V). Source: Ref 7 The H11 die steels have very high yield strengths (Fig. 8 ) and are primarily used in aircraft and missiles when high strength-to-weight ratios are desired. The H11 die steels are basically medium-carbon, 5% Cr steels with molybdenum and vanadium added. This composition air hardens from the austenitizing temperature and is tempered to a tensile strength of 1500 to 2200 MPa (215 to 320 ksi) (45 to 58 HRC). These properties apply to thin sheet as well as heavy forgings because the hardenability is fairly constant up to 38 mm (24=16in.) on a standard end-quenched specimen. Exposure of this steel at a temperature as close as 30 °C (50 °F) below the tempering temperature for 100 h or longer will have little effect on hardness and tensile strength. The high tempering temperature eliminates most residual stress. The retention of 70 to 80% of the room-temperature strength up to 540 °C (1000 °F) gives H11 steel a high strength-to-weight ratio at elevated temperatures. Additional information on H11 die steels is contained in the article "Ultrahigh-Strength Steels" in this Volume.

Mechanical Properties at Elevated Temperatures The allowable design stresses for steels at elevated temperatures may be controlled by different mechanical properties, depending on the application and temperature exposure. For applications with temperatures below the creep-temperature range, tensile strength or the yield strength at the expected service temperature generally controls allowable stresses. For temperatures in the creep range, allowable stresses are determined from either creep-rupture properties or the degree of deformation from creep. In recent years, the worldwide interest in life extension of high-temperature components has also promoted considerably more interest in elevated-temperature fatigue. This effort has led to tests and methods for evaluating the effects of creep-fatigue interaction on the life of elevated-temperature components. Ductility and toughness may also be important considerations, although ductility and toughness considerations commonly do not enter directly into the setting of allowable stresses. In elevated-temperature applications, ductility and toughness may not


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remain fixed in magnitude or character and often change with temperature and with time at temperature. The changes, which may be beneficial but are often deleterious, are of interest both at service temperature and, because of shutdowns, at ambient temperatures. Ductility is also an important factor that influences notch sensitivity and creep-fatigue interaction. The types of tests used to evaluate the mechanical properties of steels at elevated temperatures include: • • • • • •

Short-term elevated-temperature tests Long-term elevated-temperature tests Fatigue tests (including thermal fatigue and thermal shock tests) Time-dependent fatigue tests Ductility and toughness tests Short-term and long-term tests following long-term exposure to elevated temperatures

Several methods are used to interpret, interpolate, and extrapolate the data from some of these tests, as described in the section "Methods for Correlating, Interpolating, and Extrapolating Elevated-Temperature Mechanical Property Data" in this article. Short-Term Elevated-Temperature Tests Short-term elevated-temperature tests include the elevated-temperature tensile tests (described in ASTM E 21), a test for elastic modulus (ASTM E 231), compression tests, pin bearing load tests, and the hot hardness test. The mechanical properties determined by means of the tensile test include ultimate tensile strength, yield strength, percent elongation, and percent reduction in area. Because elevated-temperature tensile properties are sensitive to strain rate, these tests are conducted at carefully controlled strain rates. Tensile strength data obtained on specimens of annealed 21=4Cr-1Mo steel at various temperatures and at strain rates ranging from 2.7 × 10−6 s−1 to 144 s−1 are shown in Fig. 11 . Fig. 11 Effect of test temperature and strain rate on the strength of annealed 21=4Cr-1Mo steel. Tensile strength (a) and yield strength (b) of 21=4Cr-1Mo steel tested at various temperatures and strain rates. Source: Ref 22


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In designing components that are to be produced from low-alloy steels and to be exposed to temperatures up to 370 °C (700 °F), the yield and ultimate strengths at the maximum service temperature can be used much as they would be used in the design of components for service at room temperature. Figure 8 compares the short-time elevated-temperature yield and tensile strengths of selected alloys. Certain codes require that appropriate factors be applied in calculating allowable stresses. Elevated-temperature values of elastic modulus can be determined during tensile testing or dynamic testing by measuring the natural frequency of a test bar at the designated test temperature. Figure 12 shows values of elastic modulus at temperatures between room temperature and 650 °C (1200 °F) for several low-alloy steels, determined during static tensile loading and dynamic loading. Fig. 12 Effect of test temperature on elastic modulus, shear modulus, and Poisson's ratio. (a) Effect of test temperature on elastic modulus for several steels commonly used at elevated temperatures. Dynamic measurements of elastic modulus were made by determining the natural frequencies of test specimens; static measurements were made during tensile testing. (b) Effect of test temperature on shear modulus of 21=4Cr-1Mo steel. (c) Effect of test temperature on Poisson's ratio of 21=4Cr-1Mo steel. Source: Ref 23


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Compression tests and pin bearing load tests (ASTM E 209 and E 238) can be used to evaluate materials for applications in which the components will be subjected to these types of loading at elevated temperatures. Hot hardness tests can be used to evaluate materials for elevated-temperature service and can be applied to the qualification of materials in the same way in which room-temperature hardness tests are applied. Components for many elevated-temperature applications are joined by welding. Elevated-temperature properties of both the weld metal and the heat-affected zones can be determined by the same methods used to evaluate the properties of the base metal. Long-Term Elevated-Temperature Tests Long-term elevated-temperature tests are used to evaluate the effects of creep, which is defined as the time-dependent strain that occurs under constant load at elevated temperatures. Creep is observed in steels at temperatures above about 370 °C (700 °F). In general, creep occurs at a temperature slightly above the recrystallization temperature of a metal or alloy; at such a temperature, atoms become sufficiently mobile to allow time-dependent rearrangement of the structure. In time, creep may lead to excessive deformation and even fracture at stresses considerably below those determined in room-temperature and elevated-temperature short-term tension tests. Typical creep behavior consists of three distinct stages, as shown in Fig. 13 . Following initial elastic-plastic strain resulting from the immediate effects of the applied load, there is a region of increasing plastic strain at a decreasing strain rate (first-stage, or primary, creep). Following the primary creep region, there is a region where the creep strain increases at a minimum, and almost constant, rate of plastic strain (second-stage, or secondary creep). This nominally constant creep rate is generally known as the minimum creep rate and is widely employed in research and engineering studies. Finally, there is a region of drastically increased strain rate with rapid extension to fracture (third-stage, or tertiary creep). Tertiary creep has no distinct beginning but


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does refer to the region with an increasing rate of extension that is followed by fracture. Under certain conditions, some metals may not exhibit all three stages of plastic extension. For example, at high stresses or temperatures, the absence of primary creep is not uncommon, with secondary creep or, in extreme cases, tertiary creep following immediately upon loading. Fig. 13 Schematic representation of classical creep behavior

Of all the parameters pertaining to the creep curve, the most important for engineering applications are the creep rate and the time to rupture. These parameters are determined from long-term elevated-temperature tests that include creep, creep-rupture, and stress-rupture tests (ASTM E 139) and notched-bar rupture tests (ASTM E 292). In addition, relaxation tests (ASTM E 328) are used to evaluate the effect of creep behavior on the performance of high-temperature bolt steels. These tests are described in the article "Creep, Stress-Rupture, and Stress-Relaxation Testing" in Mechanical Testing, Volume 8 of ASM Handbook, formerly 9th Edition Metals Handbook. Creep Strength. When the rate or degree of deformation is the limiting factor, the design stress is based on the minimum (secondary) creep rate and design life after allowing for initial transient creep. The stress that produces a specified minimum creep rate of an alloy or a specified amount of creep deformation in a given time (for example, 1% total creep in 100,000 h) is referred to as the liming creep strength or limiting stress. Typical creep strengths of various low-alloy steels are shown in Fig. 3 . Table 2 also lists some suggested maximum service temperatures of various low-alloy steels based on creep rate. Figure 14 shows the 0.01%/1000 h creep strength of carbon steel as a function of room temperature tensile strength. Fig. 14 Relationship between creep strength (0.01%/1000 h) and ultimate tensile strength of a carbon steel. Creep strength estimates made using isothermal lot constant. Source: Ref 24

Stress Rupture. When fracture is a limiting factor, stress-rupture values are used in design. Stress-rupture values of various low-alloy chromium-molybdenum steels are shown in Fig. 4 , 5 , and 6 . Figures 15 and 16 show typical creep-rupture values of carbon and 1Cr-1Mo-0.25V steel, respectively. Fig. 15 Predicted 105-h creep-rupture strengths of carbon steel with (a) coarse-grain deoxidation practice and (b) fine-grain deoxidation practice. Source: Ref 24


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Fig. 16 Time-temperature-rupture data of a 1Cr-1Mo-0.25V steel

It should be recognized that long-term creep and stress-rupture values (for example, 100,000 h) are often extrapolated from shorter-term tests. Whether these property values are extrapolated or determined directly often has little bearing on the operating life of high-temperature parts. The actual material behavior is often difficult to predict accurately because of the complexity of the service stresses relative to the idealized, uniaxial loading conditions in the standardized tests and because of the attenuating factors such as cyclic loading, temperature fluctuations, or metal loss from corrosion. For those alloys in which failure occurs before a well-defined start of tertiary creep, it is useful to use notched specimens or specimens with both smooth and notched test sections (with the cross-sectional area of the notch equal to that of the smooth test section). If the material is notch sensitive, the specimen will fail in the notch before failure occurs in the smooth section. It has been well recognized for many years that notch sensitivity is related to creep ductility. It has been suggested that a minimum smooth-bar creep ductility of about 10% in terms of reduction in area may be desirable for avoidance of notch sensitivity (Ref 25 and 26). Limited published data on notched stress-rupture properties of low-alloy ferritic steels for elevated temperatures indicate that these steels generally are not notch sensitive. Representative stress-rupture data for notched and unnotched specimens of AISI 603 steel are presented in Fig. 17 . Fig. 17 Effect of notch on stress-rupture behavior. Stress-rupture behavior of smooth (K = 1.0) and notched specimens of AISI 603 steel tested at 595 °C (1100 °F). All specimens were normalized at 980 °C (1800 °F) and tempered 6 h at 675 °C (1250 °F). Source: Ref 27


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Relaxation Tests. Creep tests on metals are usually carried out by keeping either the applied load or the stress constant and noting the specimen strain as a function of time. In another type of test, known as the stress relaxation test, a sample is first deformed to a given strain and then the stress is measured as a function of time such that the total strain remains constant. Stress relaxation tests are more difficult to carry out than ordinary creep tests and are more difficult to interpret. However, stress relaxation is an important elevated-temperature property in the design of bolts or other devices intended to hold components in contact under pressure. If the service temperature is high enough, the extended-time stress on the bolt causes a minute amount of creep, which results in a reduction in the restraining force. Because of their low relaxed stresses, carbon steels are usually used only at temperatures below 370 °C (700 °F). Various low-alloy steels have been widely used up to metal temperatures of about 540 °C (1000 °F). Modified 12% Cr steels can be used for slightly higher temperatures. The common austenitic stainless steels are seldom used because of their low yield strength in the annealed condition, but are used in the cold-worked condition. The superstrength alloys are usually employed only at the highest temperatures. The comparative 1000-h relaxation strengths of these classes of alloys are shown in Fig. 18 (a). More recent data are provided in Ref 29. Fig. 18 Comparison of relaxation strengths (residual stress) of various steels. (a) Comparison of low-alloy steels with superstrength alloys. (b) Low-alloy steels at 1000 h. (c) Low-alloy steels at 10,000 h. Source: Ref 28


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Carbon steel is not recognized as a high-temperature bolting material under ASTM standards or by the ASME Boiler Code. One of the most widely used low-alloy steels for moderately high-temperature bolting applications is quenched and tempered 4140, in accordance with ASTM A 193, grade B7. Its relaxation behavior is approximately indicated by the solid lines for 0.65-1.10Cr-0.10-0.30Mo steels in Fig. 18 (b) and 18 (c). The relaxation strength of 4140 is greater after normalizing and tempering than in the quenched and tempered condition. However, this steel is nearly always used in the quenched and tempered condition in order to obtain more consistent mechanical properties. Chromium-molybdenum steels similar to 4140 except that they contain approximately 0.50% Mo (A 193, grade B7A) have also been widely used. They have slightly higher relaxation strength than 4140 but are less readily available. The strongest low-alloy steels are those with approximately 1% Cr, 0.5% Mo, and 0.25% V, in the normalized and tempered condition (A 193, grade B14) or the quenched and tempered condition (A 193, grade B16). Some of these grades are produced with rather high silicon contents (~0.75%), which seems to increase resistance to tempering. These grades have been satisfactory in service up to 540 °C (1000 °F) metal temperature in the absence of excessive follow-up or retightening. However, they are somewhat notch sensitive in creep rupture and in impact at room temperature, especially in the normalized and tempered condition.


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Fatigue At room temperature and in nonaggressive environments (and except at very high frequencies), the frequency at which loads are applied has little effect on the fatigue strength of most metals. The effects of frequency, however, become much greater as the temperature increases or as the presence or corrosion becomes more significant. At high temperatures, creep becomes more of a factor, and the fatigue strength seems to depend on the total time stress is applied rather than solely on the number of cycles. The behavior occurs because the continuous deformation (creep) under load at high temperatures affects the propagation of fatigue cracks. This effect is referred to as creep-fatigue interaction. The quantification of creep-fatigue interaction effects and the application of this information to life prediction procedures constitute the primary objective in time-dependent fatigue tests. Time-dependent fatigue tests are also used to assess the effect of load frequency on corrosion fatigue. Effect of Load Frequency on Corrosion Fatigue. In aggressive environments, fatigue strength is strongly dependent on frequency. Corrosion fatigue strength (endurance limit at a prescribed number of cycles) will generally decrease as the cyclic frequency is decreased. This effect is most important at frequencies of less than 10 Hz. The frequency dependence of corrosion fatigue is thought to result from the fact that the interaction of a material and its environment is essentially a rate-controlled process. Low frequencies, especially at low strain amplitudes or when there is substantial elapsed time between changes in stress levels, allow time for interaction between material and environment; high frequencies do not, particularly when high strain amplitude is also involved. At very high frequencies or in the plastic-strain range, localized heating may seriously affect the properties of the part. Such effects normally are not considered to be related to a corrosion fatigue phenomenon. When environments have a deleterious effect on fatigue behavior, a critical range of frequencies of loading may exist in which the mechanical/environmental interaction is significant. Above this range the effect usually disappears, while below this range the effect may diminish. Creep-fatigue interaction is an elevated-temperature phenomenon that can seriously reduce fatigue life and creep-rupture strength. Figure 19 illustrates the effect of time-dependent fatigue when the elevated temperature is within the creep range of a material. Figure 19 shows a continuous strain cycling waveform (Fig. 19 a) and a hold cycling waveform (Fig. 19 b) for fatigue strength testing. Figure 19 c) shows the fatigue life from a continuous strain cycle and from cycling with two different hold times. This decrease in fatigue life with increasing hold time or decreasing frequency, which occurs at temperatures within the creep range, is referred to as time-dependent fatigue or creep-fatigue interaction. It has been attributed to a number of factors, including the formation of intergranular voids or classical creep damage (which permits intergranular crack propagation under cyclic loading conditions), environmental interaction (corrosion fatigue), mean stress effects, and microstructural instabilities of defects produced as a result of stress and/or thermal aging, irradiation damage, and fabrication processing. Fig. 19 Range of conditions to be considered in studies of elevated-temperature fatigue and the effect of continuous cycling (a) and strain hold (b) on elevated-temperature fatigue (c). Source: Ref 30


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Most of these changes can occur at elevated temperatures and are time and possibly waveform dependent. There is also ample evidence to show that rupture ductility has a major influence on creepfatigue interaction. Because this effect is believed to be caused by the influence of rupture ductility on the creep-fracture component, endurance in continuous-cycle and in high-frequency or short-hold-time fatigue tests (where fracture is fatigue-dominated) will be relatively unaffected. As the frequency is decreased or as the hold time is increased, the effect of rupture ductility becomes more pronounced. Endurance data for several ferritic steels, in relation to the range of rupture ductility exhibited by them, are illustrated in Fig. 20 . The lower the ductility, the lower the creep-fatigue endurance. In addition, long hold periods, small strain ranges, and low ductility favor creep-dominated failures, whereas short hold periods, intermediate strain ranges, and high creep ductility favor creep-fatigue-interaction failures. Similar results have been presented for austenitic stainless steels (see the article "Elevated-Temperature Properties of Stainless Steels" in this Volume). Fig. 20 Effect of ductility on endurance of ferritic steels. Source: Ref 31


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To determine the effect of cyclic loading superimposed on a constant load at elevated temperatures, several types of fatigue testing can be employed: continuous alternating stress, continuous alternating strain, tension-tension loading with the stress ratios greater than 0, and special waveforms that provide specific holding times at maximum load. Results of these tests show which factors are most contributory to deformation and fracture of the specimens for the testing conditions employed. Further information on time-dependent fatigue is available in the article "Creep-Fatigue Interaction" in Mechanical Testing, Volume 8 of ASM Handbook, formerly 9th Edition Metals Handbook. Fatigue-Crack Growth. Although the S-N curves have been used in the past as the basic tool for design against fatigue, their limitations have become increasingly obvious. One of the more serious limitations is the fact that they do not distinguish between crack initiation and crack propagation. Particularly in the low-stress regions, a large fraction of a component's life may be spent in crack propagation, thus allowing for crack tolerance over a large portion of the life. Engineering structures often contain flaws or cracklike defects that may altogether eliminate the crack-imitation step. A methodology that quantitatively describes crack growth as a function of the loading variables is, therefore, of great value in design and in assessing the remaining lives of components. Because fatigue-crack growth rates are obtained at various ∆K ranges and temperature ranges,pit is difficult to compare the p various types of materials directly. At a constant ∆K (arbitrarily chosen as 30 MPa m, or 27 ksi in:), a clear trend of crack-growth-rate increase with increasing temperature can be seen as shown in Fig. 21 . In this figure it can be seen that at temperatures up to about 50% of the melting point (550 to 600 °C, or 1020 to 1110 °F), the growth rates are relatively insensitive to temperature, but the sensitivity increases rapidly at higher temperatures. The crack-growth rates for all the materials at temperatures up to 600 °C (1110 °F) relative to the room-temperature rates can be estimated by a maximum correlation factor of 5 (2 for ferritic steels). p p Fig. 21 Variation of fatigue-crack growth rates as a function of temperature at ∆K = 30 MPa m (27 ksi in:). Source: Ref 32


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Thermal fatigue refers to the gradual deterioration and eventual cracking of a material from cyclic thermal transients. In the past, thermal fatigue traditionally has been treated as synonymous with isothermal low-cycle fatigue at the maximum temperature of the thermal cycle. Life-prediction techniques also have evolved from the low-cycle-fatigue (LCF) literature. More recently, advances in finite-element analysis and in servohydraulic test systems have made it possible to analyze complex thermal cycles and to conduct thermomechanical fatigue (TMF) tests under controlled conditions. The assumed equivalence of isothermal LCF tests and TMF tests has been brought into question as a result of a number of studies. It has been shown that for the same total strain range, the TMF test can be more damaging under certain conditions than the pure LCF test. Information on the thermal fatigue of materials is provided in Ref 33. High-cycle thermal fatigue frequently results from intermittent wetting of a hot surface by a coolant having a considerably lower temperature. In this case, thermal fatigue cracks may initiate at the surface after a sufficient number of cycles. In other cases, the thermal cycling or ratcheting may result in plastic deformation. Thermal ratcheting is the progressive cyclic inelastic deformation that occurs as a result of cyclic strains caused by thermal or secondary mechanical stresses; sustained primary loading often contributes to thermal ratcheting. Salt pots used to contain heat-treating salt are subject to thermal ratcheting whenever the salt goes through a freeze-melt cycle. Ductility and Toughness Although steels typically have adequate ambient temperature toughness and excellent elevated-temperature ductility, several embrittling machanisms can occur during elevated-temperature exposure. Consequently, ductility and toughness tests are useful in assessing embrittling mechanisms. Information on the toughness of steels is provided in the article "Notch Toughness of Steels" in this Volume. Figure 22 shows that toughness may actually decrease if steels are tempered in the range of 260 to 370 °C (500 to 700 °F). This decrease in toughness is referred to as tempered martensite embrittlement, 350 °C embrittlement, or 500 °F embrittlement and is discussed in more detail in the article "Embrittlement of Steels" in this Volume. As a result of this embrittlement, the tempering range between 260 and 370 °C (500 and 700 °F) is generally avoided in commercial practice. Another type of embrittlement⎯temper embrittlement⎯may occur in certain alloy steels as a result of holding on slow cooling through certain temperature ranges (see the text below discussing Fig. 23 ). Fig. 22 Impact toughness as a function of tempering temperature of hardened, low-alloy, medium-carbon steels. Source: Ref 34

Fig. 23 Shift in impact transition curve to higher temperature as a result of temper embrittlement produced in SAE 3140 steel by isothermal holding and furnace cooling through the critical range. Source: Ref 34


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Another method of assessing toughness is to estimate the ductile-to-brittle transition temperature by performing notched-specimen impact tests at various temperatures. Steels are susceptible to a lowering of absorbed impact energy with decreasing temperature of use or testing. This change in energy value is accompanied by a transition from a fibrous to a crystalline-appearing fracture. The temperature at which some specified level of energy absorption or fracture appearance occurs is defined as a transition temperature. Transition temperature is an important concept because it defines a change in the mode of fracture from one that is caused predominantly by a shear mechanism to one that propagates primarily by cleavage (or along the grain boundaries in the case of temper embrittlement). Shifts in the ductile-to-brittle transition temperature are measured to detect the presence of temper embrittlement, as shown in Fig. 23 . In this case, 3140 steel (containing nominally 1.15% Ni and 0.65% Cr) was embrittled by both isothermal tempering and slow furnace cooling through the critical temperature range of about 375 to 575 °C (706 to 1070 °F). Additional information on temper embrittlement is available in the article "Embrittlement of Steels" in this Volume. A third method of assessing the effects of embrittlement mechanisms is by ductility (reduction of area) measurements. Creep embrittlement effects, for example, are usually reported in terms of a ductility minimum in stress-rupture tests, while temper embrittlement is usually recorded as an upward shift in Charpy V-notch transition temperature. Creep embrittlement occurs in roughly the same temperature range as temper embrittlement, but is not reversible with heat treatment. Creep embrittlement also seems to depend on tempering reactions inside grains and on the presence of a carbide denuded zone at prior austenite grain boundaries, while segregation effects producing temper embrittlement occur at distances only a few atomic diameters from the grain boundary. Some investigators maintain that impurities known to produce temper embrittlement also contribute to the development of creep embrittlement. Some general characteristics of creep embrittlement are: • Creep embrittlement has been shown to occur in the temperature range 425 to 595 °C (800 to 1100 °F) for alloy steels having ferrite plus carbide microstructures • Creep embrittlement appears after longer times and becomes more severe the lower the position in the embrittling temperature range • Creep embrittlement is manifested by a loss and then partial recovery of stress-rupture ductility with decreasing stress • The development of embrittlement is invariably associated with a transition from transgranular to intergranular fracture. Voids and microcracks are found throughout a creep embrittled microstructure. These voids form along prior austenite grain boundaries transverse to the tensile direction • The mechanism for creep embrittlement appears to be closely associated with tempering reactions inside grains and at the grain boundaries during the creep process. The formation of fine, needlelike precipitates in grain interiors, accompanied by the development of a denuded zone and elongated alloy carbides at grain boundaries seems to contribute significantly to the


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embrittlement process • Loss in toughness produced by creep embrittlement is largely unaffected by subsequent heat treatments; and void formation caused by creep is irreversible

Long-Term Exposure Long-term exposure to elevated temperature may affect either short-term or long-term properties. For example, the initial microstructure of creep-resistant chromium-molybdenum steels consists of bainite and ferrite containing Fe3C carbides, ε carbides, and fine M2C carbides. Although a number of different carbides may be present, the principal carbide phase responsible for strengthening is a fine dispersion of M2C carbides, where M is essentially molybdenum. With increasing aging in service, or tempering in the laboratory, a series of transformations of the carbide phases takes place that eventually transform M2C into M6C and M23C (where the M in the latter two metal carbides is mostly chromium). Such an evolution of the carbide structure results in coarsening of the carbides, changes in the matrix composition, and an overall decrease in creep strength. The effect of exposure on the stress-rupture strength of two chromium-molybdenum steels is shown in Fig. 6 . Other metallurgical changes (such as spheroidization and graphitization) and corrosion effects may also occur during long-term exposure at elevated temperature. Therefore, tests after long-term exposure may be useful in determining the effect of these metallurgical changes on short-term or long-term properties. Data Presentation and Analysis Presentation of Tensile and Yield Strength. One method for comparing steels of different strengths is to report elevated-temperature strength as a percentage of room-temperature strength; this method is illustrated in Fig. 24 . The strength levels of the steels represented in Fig. 24 varied from 480 to 1100 MPa (70 to 160 ksi). Fig. 24 Ratios of elevated-temperature strength to room-temperature strength for hardened and tempered 21=4Cr-1Mo steel tempered to room-temperature tensile strengths ranging from 480 to 1100 MPa (70 to 160 ksi). (a) Tensile strength. (b) Yield strength. Source: Ref 2

Presentation of Creep Data. Four different presentations of the same creep data for 21=4Cr-1Mo steel given in Fig. 25 . In Fig. 15 (a) to (c), only the creep strain is plotted. In the isochronous stress-strain diagram (Fig. 25 d), total strain is used. The overall format of Fig. 25 (d) is particularly useful in design problems in which total strain is a major consideration. Fig. 25 Analysis of creep data. Creep behavior of 21=4Cr-1Mo steel tested at 540 °C (1000 °F). (a) Creep strain-time plot; constant-stress lines have been drawn parallel. (b) Stress-creep strain plot. (c) Stress-time plot; constant-strain lines have been drawn parallel. (d) Isochronous stress-strain curves. Source: Ref 35


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Methods for Correlating, Interpolating, and Extrapolating Elevated-Temperature Mechanical Property Data. The behavior of steels at elevated temperatures can be affected by many variables, including time, temperature, stress, and environment. A variety of methods have been devised for correlating, interpolating, and extrapolating elevated-temperature mechanical property data. Further information on the analysis of elevated-temperature data is contained in Ref 32; Mechanical Testing, Volume 8 of ASM Handbook, formerly 9th Edition Metals Handbook; and MPC-7 of the American Society of Mechanical Engineers. Larson-Miller Parameter. Several parameters have been used for comparison of and interpolation between stress-rupture data. The most widely used is the Larson-Miller parameter, P, defined by the equation: P=T (C + log t) × 10−3 (Eq 1a) where T is the test temperature in degrees Rankine, t is the rupture time in hours, and C is a constant whose value is approximately 20 for low-alloy steels. If T is given in Kelvins, the equation is: P=1.8 T (C + log t) × 10−3 (Eq 1b) The Larson-Miller parameter is used with an experimentally determined graph such as that shown in Fig. 26 to correlate stress, temperature, and rupture time. Each graph should include the ranges of time and temperature for which the data apply; extrapolation beyond these ranges is generally not appropriate. Fig. 26 Larson-Miller plot of stress-rupture behavior of 21=4Cr-1Mo steel. Variation in Larson-Miller parameter with stress to rupture for normalized and tempered and hardened and tempered specimens of 21=4Cr-1Mo steel tested between 425 and 650 °C (800 and 1200 °F) for rupture life to 10,000h; the data are grouped according to the room-temperature tensile strength of the steel. Larson-Miller plot for annealed steel included for comparison. Source: Ref 2


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A similar parameter was used by Smith (Ref 36) to describe the creep behavior of 9Cr-1Mo steel. The creep rate parameter, P′, is given by: P′=T (20−log r) × 10−3 (Eq 2a) where T is the test temperature in degrees Rankine, and r is the minimum creep rate in percent per hour. If T is given in Kelvins, the equation is: P′=1.8 T (20−log r) × 10−3 (Eq 2b) The creep parameter is used with an experimentally determined graph such as the one shown in Fig. 27 for 2.25Cr-1Mo steel. Fig. 27 Modified Larson-Miller plot of creep behavior of 21=4Cr-1Mo steel. Variation in creep rate parameter with creep stress for normalized and tempered and hardened and tempered specimens of 21=4Cr-1Mo steel tested between 425 and 650 °C (800 and 1200 °F) for test duration to 10,000 h; the data are grouped according to the room-temperature tensile strength of the steel. Creep rate data for annealed steel included for comparison. Source: Ref 2, 36

Extrapolation of Creep and Rupture Data. It should be recognized that long-term creep and stress-rupture values (for example, 100,000 h) are often extrapolated from shorter-term tests. Whether these property values are extrapolated or determined directly often has little bearing on the operating life of high-temperature parts. The actual material behaviour is often difficult to


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predict accurately because of the complexity of the service stresses relative to the idealized, uniaxial loading conditions in the standardized tests and because of the attenuating factors such as cyclic loading, temperature fluctuations, or metal loss from corrosion. Marked changes in the slope of stress-rupture curves (see, for example, the lower plot in Fig. 15 (b) near 480 °C, or 900 °F) must also be considered in data extrapolation. These changes in slope are often indicative of microstructural changes. Marked differences in slope between curves representing temperatures separated by less than 100 °C (180 °F) should be regarded as evidence that the slope of the lower-temperature curve will change over the time period of extrapolation, indicating the need for longer tests or careful approximations of the probable influence of the change in slope. Such changes in slope are almost always in the direction of lower stress-rupture strength than would be predicted by straight-line extrapolation. Because of microstructural instabilities, deviations from the ideal creep curve must also be considered. Primary creep may be virtually absent or may be excessive and extend over long periods of time. Secondary creep may persist only for very short time periods or may exhibit nonclassical behavior. The creep behavior of annealed 2.25Cr-1Mo steel, for example, exhibits creep curves that differ from a classical three-stage creep curve in that two steady-state stages occur (Ref 37). During the first steady-state stage, the creep rate is controlled by the motion of dislocations that contain atmospheres of carbon and molybdenum atom clusters, a process termed interaction solid-solution hardening (see the section "Strengthening Mechanisms" in this article). Eventually, the precipitation of Mo2C removes molybdenum and carbon from solution, and the creep rate increases to a new steady state where the creep rate is controlled by atmosphere-free dislocations moving through a precipitate field. These nonclassical curves occur at intermediate stresses. As the stress decreases, the first steady-state stage disappears because the dislocation velocity decreases and the molybdenum-carbon atmosphere will be able to diffuse with the dislocations. At high stresses, a classical curve occurs when the creep rate is controlled by a combination of processes that operate in the two steady-state stages of the nonclassical curves (Ref 37). Such factors indicate the need to experimentally check values of deformation predicted by extrapolation of secondary creep data. The extrapolation of stress-rupture ductility with parametric techniques has been considered a potential method for predicting long-term ductility from short-term tests (Ref 38) Because the stress-rupture ductility of many alloys used at elevated temperatures varies with temperature and stress, the objective is to develop a combined (stress, temperature) parameter that can be correlated to rupture ductility over a wide range of stresses and temperatures. Reference 38 compares the correlation between some parametric models and rupture ductility data for a 11=4Cr-1=2Mo steel in the temperature range of 510 to 620 °C (950 to 1150 °F). Methods for Predicting Time-Dependent Fatigue (Creep-Fatigue Interaction) Behavior. Many methods have been employed to extrapolate available data to estimate the time-dependent fatigue life of materials. Development of a mathematical formulation for life prediction is one of the most challenging aspects of creep-fatigue interaction. It is complicated by the fact that any proposed formulation must account for strain rate, relaxation at constant strain, creep at constant load, the difference between tension and compression creep and/or relaxation, or combinations of all of these. Linear damage summation is perhaps the most widely known and simplest of the many life prediction methods and has been used extensively in the evaluation of creep-fatigue interaction. It is based on a simple relation that fatigue damage can be expressed as a cycle fraction of damage and that creep damage can be expressed as a time fraction of damage. It is also assumed that these quantities can be added linearly to represent damage accumulation. Failure occurs when this summation reaches a certain value. Other methods include the ductility exhaustion approach, the frequency modified approach, and strain range partitioning. These methods are reviewed in Ref 30, 32 and in Mechanical Testing, Volume 8 of ASM Handbook, formerly 9th Edition Metals Handbook.

Corrosion Although the mechanical properties establish the allowable design-stress levels, corrosion effects at elevated temperatures often set the maximum allowable service temperature of an alloy. The following sections describe the three common forms of corrosion⎯oxidation, sulfidation, and hydrogen attack⎯that occur at elevated temperature. Corrosion considerations with liquid-metal environments are also summarized. More detailed information on corrosion and its prevention is available in Corrosion, Volume 13 of ASM Handbook, formerly 9th Edition Metals Handbook. Oxidation from steam or air is a serious problem that can occur at elevated temperatures. When metal is exposed to an oxidizing gas at elevated temperature, corrosion can occur by direct reaction with the gas. This type of corrosion is referred to as tarnishing, high-temperature oxidation, or scaling. The rate of attack increases substantially with temperature. The surface film typically thickens as a result of reaction at the scale/gas or metal/scale interface due to cation or anion transport through the scale, which behaves as a solid electrolyte. Alloys intended for high-temperature applications are designed to have the capability of forming protective oxide scales. Chromium provides oxidation resistance in alloy steels, and Fig. 28 compares the loss by scaling for alloys with varying levels of chromium. Silicon can also improve oxidation resistance (Fig. 28 ), although it also reduces creep strength and may promote temper embrittlement when other impurities are present. Fig. 28 Effect of temperature on metal loss from scaling for several carbon and alloy steels in air


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The scaling data given in Fig. 28 were obtained in the presence of air. If other variables affecting oxidation are changed, such as gas composition, the heating method, temperature, pressure, or velocity, different rates of scaling can be expected. Elements such as sulfur, vanadium, and sodium can change the nature of metal oxidation, sometimes increasing it to a catastrophic level of several inches per year. At elevated temperatures, steam decomposes at metal surfaces to hydrogen and oxygen and may cause steam oxidation of steel, which is somewhat more severe than air oxidation at the same temperature. Fluctuating steam temperatures tend to increase the rate of oxidation by causing scale to spall, which exposes fresh metal to further attack. Table 1 gives the maximum-use temperatures for several boiler alloys for which code standards exist. The strength criteria are based on the wall midsection temperatures, which are typically 25 °C (45 °F) lower than the outer surface temperature. In a water environment, corrosion is significantly influenced by the concentrations of dissolved species, pH, temperature, suspended particles, and bacteria. Temperature plays a dual role with respect to oxygen corrosion. Increasing the temperature will reduce oxygen solubility. In open systems, in which oxygen can be released from the system, corrosion will increase up to a maximum at 80 °C (175 °F), where the oxygen solubility is 3 mg/L. Beyond this temperature, the reduced oxygen content limits the oxygen reduction reaction, preventing occurrence of the iron dissolution process. Thus, the corrosion rate of carbon steel decreases, and at boiling water conditions, the temperature effect is similar to room temperature with a high oxygen content. For closed systems, in which oxygen cannot escape, corrosion continues to increase linearly with temperature. The other factors affected by temperature are the diffusion of oxygen to the metal surface, the viscosity of water, and solution conductivity. Increasing the temperature will increase the rate of oxygen diffusion to the metal surface, thus increasing corrosion rate because more oxygen is available for the cathodic reduction process. The viscosity will decrease with increasing temperature, which will aid oxygen diffusion. Sulfidation. Corrosion by various sulfur compounds at temperatures between 260 and 540 °C (500 and 1000 °F) is a common problem in many petroleum-refining processes and, occasionally, in petrochemical processes. When the sulfur activity (partial pressure and concentration) of the gaseous environment is sufficiently high, sulfide phases, instead of oxide phases, can be formed. In the majority of environments encountered in practice by oxidation-resistant alloys, Al2O3 or Cr2O3 should form in preference to any sulfides, and destructive sulfidation attack occurs mainly at sites where the protective oxide has broken down. The role of sulfur, once it has entered the alloy, appears to be to tie up the chromium and aluminum as sulfides, effectively redistributing the protective scale-forming elements near the alloy surface and thus interfering with the process of formation or re-formation of the protective scale. If sufficient sulfur enters the alloy so that all immediately available chromium or aluminum is converted to sulfides, then the less stable sulfides of the base metal may form because of morphological and kinetic reasons. It is these base metal sulfides that are often responsible for the observed accelerated attack, because they grow much faster than the oxides or sulfides of chromium or aluminum; in addition, they have relatively low melting points, so that molten slag phases are often possible. Sulfur can also transport across continuous protective scales of Al2O3 and Cr2O3 under certain conditions, with the result that discrete sulfide precipitates can be observed immediately beneath the scales on alloys that are behaving in a protective manner. For the reasons indicated above, as long as the amount of sulfur present as sulfides is small, there is little danger of accelerated attack. However, once sulfides have formed in the alloy, there is a tendency for the sulfide phases to be preferentially oxidized by the encroaching reaction front and for the sulfur to be displaced inward, forming new sulfides deeper in the alloy, often in grain boundaries or at the sites of other chromium- or aluminum-rich phases, such as carbides. In this way, fingerlike protrusions of oxide/sulfide can be formed from the alloy surface inward, which may act to localize stress or otherwise reduce the load-bearing section. The relative corrosivity of sulfur compounds generally increases with temperature. Depending on the process particulars,


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corrosion is in the form of uniform thinning, localized attack, or erosion-corrosion. Corrosion control depends almost entirely on the formation of protective metal sulfide scales that exhibit parabolic growth behavior. In general, nickel and nickel-rich alloys are rapidly attacked by sulfur compounds at elevated temperatures, while chromium-containing steels provide excellent corrosion resistance (as does aluminum). The combination of hydrogen sulfide and hydrogen can be particularly corrosive, and as a rule, austenitic stainless steels are required for effective corrosion control. The effect of temperature and alloy compositions on sulfidic corrosions depends on the sulfur compounds present. Figure 29 shows the rates of sulfur corrosion of various steels as a function of temperature. These so-called McConomy curves can be used to predict the relative corrosivity of crude oils and their various fractions (Ref 40). Although this method relates corrosivity to total sulfur content, and therefore does not take into account the variable effects of different sulfur compounds, it can provide reliable corrosion trends if certain corrections are applied. Plant experience has shown that the McConomy curves, as originally published, tend to predict excessively high corrosion rates. The curves apply only to liquid hydrocarbon streams containing 0.6 wt% S (unless a correction factor given in Fig. 30 for sulfur content is applied) and do not take into account the effects of vaporization and flow regime. The curves can be particularly useful, however, for predicting the effect of operational changes on known corrosion rates. Fig. 29 Modified McConomy curves showing the effect of temperature on high-temperature sulfidic corrosion of various steels and stainless steels. Source: Ref 39

Fig. 30 Effect of sulfur content on corrosion rates predicted by modified McConomy curves in the temperature range of 290 to 400 °C (550 to 750 °F). Source: Ref 39

Over the years, it has been found that corrosion rates predicted by the original McConomy curves should be decreased by a factor of roughly 2.5, resulting in the modified curves shown in Fig. 29 . The curves demonstrate the beneficial effects of alloying steel with chromium in order to reduce corrosion rates. Corrosion rates are roughly halved when the next higher grade of low-alloy steel (for example, 2.25Cr-1Mo, 5Cr-0.5Mo, 7Cr-0.5Mo, or 9Cr-1Mo steel) is selected. Essentially, no corrosion occurs with stainless steels containing 12% or more chromium. Although few data are available, plant experience has shown that corrosion rates start to decrease as temperatures exceed 455 °C (850 °F). Two explanations frequently offered for this phenomenon are the possible decomposition of reactive sulfur compounds and the formation of a protective coke layer.


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Sulfidic Corrosion With the Presence of Hydrogen. The presence of hydrogen increases the severity of high-temperature sulfidic corrosion. Hydrogen converts organic sulfur compounds to hydrogen sulfide; corrosion becomes a function of hydrogen sulfide concentration (or partial pressure). A number of researchers have proposed various corrosion rate correlations for high-temperature sulfidic corrosion in the presence of hydrogen (Ref 41, 42, 43, 44, 45, 46, and 47), but the most practical correlations seem to be the so-called Couper-Gorman curves. The Couper-Gorman curves are based on a survey conducted by National Association of Corrosion Engineers (NACE) Committee T-8 on Refining Industry Corrosion (Ref 48). The Couper-Gorman curves differ from those previously published in that they reflect the influence of temperature on corrosion rates throughout a whole range of hydrogen sulfide concentrations. Total pressure was found not to be a significant variable between 1 and 18 MPa (150 and 2650 psig). It was also found that essentially no corrosion occurs at low hydrogen sulfide concentrations and temperatures above 315 °C (600 °F) because the formation of iron sulfide becomes thermodynamically impossible. Curves are available for carbon steel, 5Cr-0.5Mo steel, 9Cr-1Mo steel, 12% Cr stainless steel, and 18Cr-8Ni austenitic stainless steel. For the low-alloy steels, two sets of curves apply, depending on whether the hydrocarbon stream is naphtha or gas oil. The curves again demonstrate the beneficial effects of alloying steel with chromium to reduce the corrosion rate. Modified Couper-Gorman curves are shown in Fig. 31 . To facilitate the use of these curves, original segments of the curves were extended (dashed lines). In contrast to sulfidic corrosion in the absence of hydrogen, there is often no real improvement in corrosion resistance unless chromium content exceeds 5%. Therefore, the curves for 5Cr-0.5Mo steel also apply to carbon steel and low-alloy steels containing less than %5 Cr. Stainless steels containing at least 18% Cr are often required for essentially complete immunity to corrosion. Because the Couper-Gorman curves are primarily based on corrosion rate data for an all-vapor system, partial condensation can be expected to increase corrosion rates because of droplet impingement. Fig. 31 Effect of temperature and hydrogen sulfide content on high-temperature H2S/H2 corrosion of (a) carbon steel, (b) 5Cr-0.5Mo steel, and (c) 9Cr-1Mo steel. These corrosion rates are based on the use of gas oil desulfurizers; corrosion rates with naphtha desulfurizers may be slightly less severe. Source: Ref 39

When selecting steels for resistance to high-temperature sulfidic corrosion in the presence of hydrogen, the possibility of high-temperature hydrogen attack should be considered. Conceivably, this problem arises when carbon steel and low-alloy steels containing less than 1% Cr are chosen for temperatures exceeding 260 °C (500 °F) and hydrogen partial pressures above 700 kPa (100 psia) and when corrosion rates are expected to be relatively low. Hydrogen Damage. Because iron-base alloys are principal materials of construction, these alloys have been the focus of most of the studies relating to hydrogen effects. In addition, ferritic (body-centered cubic) steels have a particular sensitivity to hydrogen. For these reasons, hydrogen effects on steel are important. Such hydrogen effects have been thoroughly described in a review of hydrogen damage (Ref 49). Hydrogen damage includes hydrogen embrittlement, hydrogen attack, and hydrogen blistering. In ferrous alloys, embrittlement by hydrogen is generally restricted to those alloys having a hardness of 22 HRC or greater. The other forms of hydrogen damage, such as hydrogen attack or hydrogen blistering, are associated with low-alloy or carbon steels. Thermodynamic calculations by Odette (Ref 50) have also shown that the carbide M3C is significantly less stable in hydrogen environments than alloy carbides such as M2C, M7C3, M23C6, and M6C. As shown in Fig. 32 , this M3C carbide can persist in normalized 2.25Cr-1Mo steel for up to 50 h at 700 °C (1290 °F) and hence may be present during service. When alloy content is


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increased to 3Cr-1.5Mo, Ritchie et al. (Ref 51) have shown that the tempering kinetics are significantly accelerated and that M3C can be eliminated from the microstructure within 1 h at 700 °C (1290 °F). Hydrogen exposures of 3Cr-1.0Mo-0.5Ni steel at 600 °C (1110 °F) and 17 MPa (2.5 ksi) for 100 h gave no indication of the voids characteristic of hydrogen attack on these steels (Ref 51). Fig. 32 Isothermal diagram showing the sequence of carbide formation on tempering of normalized 21=4Cr-1Mo steel. Source: Ref 12

Hydrogen attack is a damage mechanism that is associated with carbon and low-alloy steels exposed to hydrogen-containing environments at temperatures above 220 °C (430 °F) (Ref 49). With hydrogen at elevated temperatures and pressures, there is increasing availability of atomic hydrogen that can easily penetrate metal structures and react internally with reducible species. Exposure to the environment is known to result in a direct chemical reaction with the carbon in the steel. The reaction occurs between absorbed hydrogen and the iron carbide phase, resulting in the formation of methane: 2H2 + Fe3C → CH4 + 3Fe Unlike nascent hydrogen, the resulting methane gas does not dissolve in the iron lattice. Internal gas pressures develop, leading to the formation of voids, blisters, or cracks. The generated defects lower the strength and ductility of the steel. Because the carbide phase is a reactant in the mechanism, its absence in the vicinity of generated defects serves as direct evidence of the mechanism itself. Hydrogen attacks occur in carbon steels and can lead to fissuring of the steel. Alloy steels with stable carbides, such as chromium and molybdenum carbides, are less susceptible to this form of attack. For example, 2.25Cr-1Mo suffers some decarburization in high-temperature high-pressure hydrogen, but is less likely to fissure than carbon steel. Hydrogen attack does not occur in austenitic stainless steels (Ref 49). The susceptibility of steels to attack by hydrogen can be judged from the Nelson curves, which indicate the regions of temperature and pressure in which a variety of steels will suffer attack. Nelson curves for various alloy steels are shown in Fig. 33 . In Fig. 33 (b), the Nelson curve indicates only the operating limits for the surface decarburization of 21=4Cr-1Mo steel. There are, however, indications that the limiting condition for using quenched and tempered 21=4Cr-1Mo steel will not be decarburization but rather the loss of integrity through methane bubble growth (Ref 54). This may only be true if fatigue is not a significant design consideration, because decarburization can affect fatigue strength. The top curve in Fig. 33 (b) shows an estimate of the operating limits for the formation of methane bubbles (that is, hydrogen attack) in 21=4Cr-1Mo steel (Ref 53). Fig. 33 (a) Nelson curves defining the operating limits of various alloys in a hydrogen environment. Curves adapted from the chart of Nelson. Source: Ref 52. (b) Nelson curves for three steels given in Ref 36 and an estimate of the operating limit (solid line) for the formation of methane bubbles in 2.25Cr-1Mo steel. Source: Ref 53


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Hydrogen blistering is a mechanism that involves the hydrogen damage of unhardened steel near ambient temperature. It is known that the entry of atomic hydrogen into steel can result in its collection, as the molecular species, at internal defects or interfaces. If the entry kinetics are substantial (promoted by an acidic environment, high corrosion rates, and cathodic poisons), the resulting internal pressure will cause internal separation (fissuring or blistering) of the steel. Such damage typically occurs at large, elongated inclusions and results in delaminations known as hydrogen blisters. Field experience indicates that fully killed steels are more susceptible than semikilled steels (Ref 55), but the nature and size of the original inclusions appear to be the key


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factors with regard to susceptibility. Rimmed steels or free-machining grades with high levels of sulfur or selenium would most likely show a high susceptibility to blistering. Stepwise cracking at the ends of blisters indicates an effect of elongated inclusions in the delamination process (Ref 49, 55). Similar stepwise cracking occurs in the hydrogen-induced failure of low-alloy pipeline steels (Ref 56). Both stepwise cracking and blistering appear to be limited to environments in which acidic corrosion occurs and in which cathodic poisons, such as sulfide, are present to promote hydrogen entry. Hydrogen embrittlement, unlike hydrogen attack or blistering, can occur without immediate and resolvable damage within the metal structure. In this respect, hydrogen embrittlement is a somewhat reversible process. For example, hydrogenation plant equipment, operating at about 540 °C (1000 °F) with absorbed hydrogen in the steel, is cooled from operating temperature at 30 to 40 °C (50 to 75 °F) per hour with no breakage. This cooling rate is apparently slow enough to allow most of the absorbed hydrogen to diffuse from the metal without causing excessive embrittlement. Another way of removing detrimental atomic hydrogen derives from its mobility at higher temperatures. A bake-out cycle, involving temperatures of 175 to 205 °C (350 to 400 °F), allows the diffusion and escape of hydrogen from the metal or alloy. If the hydrogen charging conditions were not severe enough to cause internal damage, the bake-out cycle (described in Federal Specification QQC-320) will restore full ductility. Susceptibility to hydrogen embrittlement is strongly influenced by the strength level of the metal or alloy. In steels, untempered martensite is the most susceptible phase. Lamellar carbide structures are less desirable than those with spheroidized structures. In general, iron-base alloys with a ferritic or martensitic structure are restricted to a maximum hardness of 22 HRC. Most other alloys are restricted to a maximum hardness of 35 HRC. There are exceptions in both cases. The procedures for materials testing in a wet hydrogen sulfide environment, which is the most aggressive in promoting hydrogen entry, are discussed in NACE TM-01-77. Further information on hydrogen embrittlement is provided in the article "Embrittlement of Steels" in this Volume. Resistance of Liquid-Metal Corrosion. The following sections describe the resistance of steels to various liquid-metal environments. Liquid metals can attack steels in various ways. One form of attack that can occur is intergranular penetration and/or dissolution by liquid metal. The unfortunate aspect of this mode of attack is that it can result in a loss of strength without any large weight loss or change in appearance. In this respect, it resembles the more familiar aqueous intergranular corrosion. Additional information on liquid-metal corrosion can be found in Corrosion, Volume 13 of ASM Handbook, formerly 9th Edition Metals Handbook. Another phenomenon is liquid-metal embrittlement, which requires the presence of both stress and a liquid metal. The cracks that occur during the embrittlement process may be intergranular or transgranular. The process seems to be similar in many ways to stress-corrosion cracking. Steels have been reported to undergo embrittlement by lithium, indium, cadmium, zinc, tellurium, and various lead-tin solders. Additional information on liquid-metal embrittlement can be found in the article "Embrittlement of Steels" in this Volume. Sodium and Sodium-Potassium Alloys. Plain carbon and low-alloy steels are generally suitable for long-term use in these media at temperatures to 450 °C (840 °F). Beyond these temperatures, stainless steels are required. The principal disadvantage of ferritic steels in sodium systems is the decarburization potential and its possible effect on the mechanical properties of the ferritic steel and the other system materials. Reference 5 considers aspects of low-carbon 21=4Cr-1Mo steel as the construction material for the sodium-heated steam generator of a liquid-metal fast breeder reactor. Lithium is somewhat more aggressive to plain carbon steels than sodium or sodium-potassium. As a result, low-alloy steels should not be considered for long-term use above 300 °C (570 °F). At higher temperatures, the ferritic stainless steels show better results. Cadmium. Low-alloy steels exhibit good serviceability to 700 °C (1290 °F). Zinc. Most engineering metals and alloys show poor resistance to molten zinc, and carbon steels are no exception. Antimony. Low-carbon steels have poor resistance to attack by antimony. Mercury. Although plain carbon steels are virtually unattacked by mercury under nonflowing or isothermal conditions, the presence of either a temperature gradient or liquid flow can lead to drastic attack. The corrosion mechanism seems to be one of dissolution, with the rate of attack increasing rapidly with temperature above 500 °C (930 °F). Alloy additions of chromium, titanium, silicon, and molybdenum, alone or in combination, show resistance to 600 °C (1110 °F). Where applicable, the attack of ferrous alloys by mercury can be reduced to negligible amounts by the addition of 10 ppm Ti to the mercury; this raises the useful range of operating temperatures to 650 °C (1200 °F). Additions of metal with a higher affinity for oxygen than titanium, such as sodium or magnesium, may be required to prevent oxidation of the titanium and loss of the inhibitive action. Aluminum. Plain carbon steels are not satisfactory for the long-term containment of molten aluminum. Gallium is one of the most aggressive of all liquid metals and cannot be contained by carbon or low-alloy steels at elevated temperatures. Indium. Carbon and low-alloy steels have poor resistance to molten indium. Lead, Bismuth, Tin, and Their Alloys. Low-alloy steels have good resistance to lead up to 600 °C (1110 °F), to bismuth up to 700 °C (1290 °F), and to tin only up to 150 °C (300 °F). The various alloys of lead, bismuth, and tin are more aggressive.

Factors Affecting Mechanical Properties The factors affecting the mechanical properties of steels include the nature of the strengthening mechanisms, the microstructure, the heat treatment, and the alloy composition. This section describes these factors, with particular emphasis on chromium-molybdenum steels (especially 21=4Cr-1Mo) used for elevated-temperature service.


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In addition, various service factors such as thermal exposure and environmental conditions can induce metallurgical changes, which may affect the mechanical properties of steels used at elevated temperatures. These metallurgical changes include spheroidization, graphitization, decarburization, and carburization. Depending on the temperature and exposure environment, ferritic steels used at elevated temperatures may also be susceptible to embrittlement phenomena such as temper embrittlement, temper martensite embrittlement, creep embrittlement, hydrogen embrittlement, and liquid-metal embrittlement. Embrittlement phenomena are discussed in the article "Embrittlement of Steels" in this Volume. Strengthening Mechanisms. The creep strength of a steel is affected by the typical strengthening mechanisms⎯namely, grain refinement, solid-solution hardening, and precipitation hardening. Of these various strengthening mechanisms, the refinement of grain size is perhaps the most unique because it is the only strengthening mechanism that also increases toughness. Figure 14 shows the effect of grain size on the creep strength of a carbon steel. The creep strength of chromium-molybdenum steels is mainly derived from a complex combination of solid-solution (primarily interaction solid-solution strengthening) and precipitation effects, as illustrated in Fig. 34 . In the early stages of creep, solid-solution effects are the largest contributor to creep resistance. As time progresses, the precipitation of carbides (primarily Mo2C in the case of molybdenum steels) contributes more to the creep resistance. As time progresses still further, the strengthening effect of the carbides is reduced as the carbides coarsen (Ostwald ripening) and diffuse into stabler but weaker structures. Both of these strengthening mechanisms become unstable at high temperatures. In solid-solution hardening, an increase in temperature increases the diffusion rates of solute atoms in the dislocation atmospheres while at the same time dispersing the atoms of the atmospheres, with both effects making it easier for dislocations to move. In precipitation hardening, heating of the alloy to an excessively high temperature can cause solutionizing of the precipitates. At intermediate temperatures, the precipitates can coarsen and become less-effective impediments to dislocation motion. High stresses and high-strain cyclic loading also can lead to accelerated softening. Fig. 34 Schematic of changes in creep strengthening contributions at 550 °C (1020 °F) in (a) normalized molybdenum steel and (b) normalized and tempered molybdenum steel. Source: Ref 57

The solid-solution strengthening effect illustrated in Fig. 34 occurs primarily from a process termed interaction solid-solution hardening (or strengthening), which is a mechanism that involves the interaction of substitutional and interstitial solutes (Ref 57, 58, and 59). This process occurs in ferritic alloys that contain in solid solution interstitial and substitutional elements that have an affinity for each other. As a result of this strong attraction, atom pairs or clusters could form dislocation atmospheres that hinder dislocation motion and therefore strengthen the steel. Other solid-solution effects from either a pure substitutional solute or a pure interstitial solute do not alone provide significant creep strengthening in carbon manganese steels and molybdenum steels. The addition of interstitial solutes to iron has no significant creep-strengthening effect above 450 °C (840 °F), while the substitutional solutes manganese, chromium, and molybdenum give rise to only modest increases in strength in the absence of interstitial solutes (Ref 59). However, when certain combinations of substitutional and interstitial solutes are present together (for example, manganese-nitrogen, molybdenum-carbon, and molybdenum-nitrogen), there is a substantial increase in creep strength (Ref 59). These combinations give rise to the strengthening process of interaction solid-solution hardening. Precipitation-strengthening effects are probably negligible in the carbon-manganese steels typically used for elevated-temperature applications (Ref 57), although strengthening by fine NbC particles has been observed (Ref 60). Precipitation strengthening is more significant in molybdenum steels, for which the strengthening precipitates are mainly Mo2C


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and Mo2N. Further increases in precipitation strengthening can be achieved with additions of niobium or vanadium to chromium-molybdenum steels. The stability of the carbides increases in the following order of alloying elements: chromium, molybdenum, vanadium, and niobium. Fine and closely dispersed precipitates of NbC and VC are thus desirable, followed by the other carbides. This precipitation strengthening effect in creep-resistant chromium-molybdenum steels is related to secondary hardening, as discussed below. Secondary Hardening. If the mechanical properties of tempered steels need to be maintained at elevated service temperatures, the problem is to reduce the amount of softening during tempering so that higher strength (hardness) can be achieved at higher temperatures. One way to reduce softening is with strong carbide formers such as chromium, molybdenum, and vanadium. These carbide formers induce an effect known as secondary hardening. Without these elements, iron-carbon alloys and low-carbon steels soften rapidly with increasing tempering temperature, as shown in Fig. 35 . This softening is largely due to the rapid coarsening of cementite with increasing tempering temperature, a process dependent on the diffusion of carbon and iron. If present in a steel in sufficient quantity, however, the carbide-forming elements not only retard softening but also form fine alloy carbides that produce a hardness increase at higher tempering temperatures. This hardness increase is frequently referred to as secondary hardening. This hardening can also occur during elevated-temperature service and is related to creep strength, as shown in Fig. 36 . Fig. 35 Decrease in hardness with increasing tempering temperature for steels of various carbon contents. Source: Ref 61

Fig. 36 Relationship between change in creep rate and change in room-temperature hardness during creep of normalized 1% Mo steel tested at 123 MPa (17.8 ksi) at 550 °C (1020 °F). Under these test conditions, secondary creep coincided with maximum precipitation hardening. Source: Ref 57


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Secondary hardening allows higher tempering temperatures, and this increases the range of service temperatures. Figure 37 shows secondary hardening in a series of steels containing molybdenum. The secondary hardening peaks develop only at high tempering temperatures because alloy carbide formation depends on the diffusion of the carbide-forming elements, a more sluggish process than that of carbon and iron diffusion. As a result, not only is a finer dispersion of particles produced but also the alloy carbides, once formed, are quite resistant to coarsening. The latter characteristic of the fine alloy carbides is used to advantage in tool steels that must not soften even though high temperatures are generated by their use in hot-working dies or high-speed machining. Also, ferritic low-carbon steels containing chromium and molybdenum are used in pressure vessels and reactors operated at temperatures around 540 °C (1000 °F) because the alloy carbides are slow to coarsen at those temperatures. Fig. 37 Retardation of softening and secondary hardening during tempering of steels with various molybdenum contents. Source: Ref 61

The beneficial property changes from secondary hardening can be improved by increasing the intensity of secondary hardening, by decreasing the rate of overaging of the secondary-hardening carbide, or by increasing the temperature of secondary hardening. The intensity of secondary hardening can be increased by increasing the mis-match between the carbide precipitate and the matrix (Ref 62). Although this tends to cause more rapid overaging, the net effect can be beneficial, so that a higher strength after tempering is achieved. Increased mis-match is produced by: • Increasing the lattice parameter of the carbide precipitate • Decreasing the lattice parameter of the matrix


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The carbide Mo2C can dissolve both chromium and vanadium. Chromium, being a smaller atom than molybdenum, reduces the lattice parameter of Mo2C, but vanadium increases it. Chromium therefore tends to decrease the intensity of secondary hardening. Moreover, chromium causes the Mo2C carbide to be less stable (that is, to give maximum secondary hardening at lower tempering temperatures) and accelerates overaging (Fig. 38 ). On the other hand, vanadium increases the lattice parameter of Mo2C and stabilizes the carbide. The result is a greater intensity of secondary hardening. Fig. 38 Effect of chromium on the tempering characteristics of 0.45C-1.75Mo-0.75V steels. Source: Ref 62

Effects of Microstructure. It is widely accepted that the strength and impact toughness of carbon and chromium-molybdenum steels with fully bainitic microstructures are better than those with a ferritic-bainitic microstructure. Bainitic microstructures also have better creep resistance under high-stress, short-time conditions, but degrade more rapidly at high temperatures than pearlitic structures. As a result, ferrite-pearlite material has better intermediate-term, low-stress creep resistance. Because both microstructures will eventually spheroidize, it is expected that over long service lives the two microstructures will converge to similar creep strengths. This convergence can be estimated to occur in about 50,000 h at 540 °C (1000 °F) for 2.25Cr-1Mo steel, based on the limited data presented in Fig. 39 . Investigations of chromium-molybdenum steels for one application concluded that tempered bainite is the optimum microstructure for creep resistance (Ref 64). However, the carbide precipitates are also an important microstructural factor in achieving optimum creep behavior, and for some microstructures an untempered condition may be desirable (see the following section "Effects of Heat Treatment" in this article). Moreover, even though bainitic microstructures improve strength, toughness, and creep resistance, chromium-molybdenum steels with bainitic and tempered martensitic microstructures also undergo strain softening during mechanical cycling. This effect of strain softening of bainitic chromium-molybdenum steels has undergone several investigations (Ref 18, 19, 20, and 21). Fig. 39 Variation in stress-rupture strength of 21=4Cr-1Mo steel under different heat treatments. QT, quenched and tempered; NT, normalized and tempered; A, annealed; UTS, ultimate tensile strength. Source: Ref 63


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Microstructure may also influence the carbide precipitation and strengthening mechanism of chromium-molybdenum steels. In 21=4Cr-1Mo steel, for example, precipitation reactions are known to occur much more rapidly in bainite than in proeutectoid ferrite (Ref 12). In addition, the interaction solid-solution strengthening of 2 1=4Cr-1Mo steel is influenced by microstructure. In tensile studies, it was concluded that interaction solid-solution hardening in bainitic (normalized and tempered) 2.25Cr-1Mo steel is due to chromium-carbon interactions, while it is due to molybdenum-carbon interactions in the proeutectoid ferrite of annealed steel (Ref 65). Effects of Heat Treatment. Figure 40 shows the general effect of three heat treatments on the creep-rupture strength of 2.25Cr-1Mo steel. Like long-term exposure (Fig. 39 ), the creep-rupture strengths converge in Fig. 40 . Fig. 40 Influence of heat treatment on 105−h creep-rupture strength of 21=4Cr-1Mo steel. Source: Ref 66

The use of tempering is also an important factor that influences the level of precipitation strengthening and solid-solution strengthening in chromium-molybdenum steel (Fig. 34 ). In a normalized molybdenum steel (Fig. 34 a), the initial contribution from solid-solution strengthening is greater than that of the normalized and tempered steel. In the normalized and tempered molybdenum steels (Fig. 34 b), the initial contribution from precipitation strengthening will be larger than that from the normalized steel. In addition, the precipitation-strengthening effect in the normalized and tempered steel will reach a maximum and begin to decline at an earlier stage due to the earlier incidence of overaging in tempered material. This is a potential consideration in applications requiring creep resistance over long times and at high temperatures. As noted in the previous section "Effects of Microstructure," an investigation for one application concluded that tempered bainite resulted in the optimum creep resistance. In ferrite-pearlite or ferrite-bainite structures, however, it has been suggested that the best creep resistance at relatively high stresses is obtained in the untempered condition, because the dislocations introduced on loading are then to nucleate a finer dispersion of particles in ferrite grains than is obtained by tempering in the absence of strain (Ref 57). Application of this concept does not apply to bainitic structures. In bainitic steels, where the dislocation density is already higher than that introduced upon straining of a ferrite/pearlite steel, the use of untempered structures is unlikely to prove beneficial to short-term creep strength (Ref 57). Ultimately, it is the balance of hardness (or strength) and toughness required in service that determines the conditions of tempering for a given application. Figure 41 shows


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the variation of properties from the tempering of a modified 9Cr-1Mo alloy (9Cr-1Mo with 0.06 to 0.10 wt% Nb and 0.18 to 0.25 wt% V). Fig. 41 Room-temperature properties of two heats (open or closed symbols) of a modified 9Cr-1Mo steel correlated with the Holloman-Jaffe (HJ) tempering parameter. (a) Hardness. (b) Charpy energy. (c) 0.2% yield strength. (d) Total elongation at room temperature. Source: Ref 7

Effects of Composition. The mechanical properties of carbon and low-alloy steels are determined primarily by composition and heat treatment. The effects of alloying elements in annealed, normalized and tempered, and quenched and tempered steels are discussed below. Carbon increases both the strength and hardenability of steel at room temperature but decreases the weldability and impact toughness. In plain carbon and carbon-molybdenum steels intended for elevated-temperature service, carbon content is usually limited to about 0.20%; in some classes of tubing for boilers, however, carbon may be as high as 0.35%. For chromium-containing steels, carbon content is usually limited to 0.15%. Carbon increases short-term tensile strength but does not add appreciably to creep resistance at temperatures above 540 °C (1000 °F) because carbides eventually become spheroidized at such temperatures. Manganese, in addition to its normal function of preventing hot shortness by forming dispersed manganese sulfide inclusions, also appears to enhance the effectiveness of nitrogen in increasing the strength of plain carbon steels at elevated temperatures. Manganese significantly improves hardenability, but contributes to temper embrittlement. Phosphorus and sulfur are considered undesirable because they reduce the elevated-temperature ductility of steel. This reduction in ductility is demonstrated by reductions in stress-rupture life and thermal fatigue life. Phosphorus contributes to temper embrittlement. Silicon increases the elevated-temperature strength of steel. It also increases the resistance to scaling of the low-chromium steels in air at elevated temperatures. Silicon is one factor in temper embrittlement. Chromium in small amounts (~0.5%) is a carbide former and stabilizer. In larger amounts (up to 9% or more), it increases the resistance of steels to corrosion. Chromium also influences hardenability. The effect of chromium in ferritic creep-resistant steels is complex. By itself, chromium gives some enhancement of creep strength, although increasing the chromium content in lower-carbon grades does not increase resistance to deformation at elevated temperatures (Ref 59). When added to molybdenum steel, chromium generally leads to some reduction in creep strength (Ref 67) such as that shown in Fig. 42 . For the 1.0Mo steel in Fig. 42 , the optimum creep strength occurs with about 2.25% Cr. Chromium is most effective in strengthening molybdenum steels (0.5 to 1.0% Mo) when it is used in amounts of 1 to 21=2%. Fig. 42 Effect of chromium on the creep strength (stress to produce a minimum creep rate of 0.0001% per hour) of several steels containing small amounts of molybdenum, silicon, and aluminum at 540 °C (1000 °F). Source: Ref 68


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Figure 43 summarizes the effects of chromium content on the tensile and yield strengths of chromium-molybdenum steels containing 0.5 to 1.0% Mo and various amounts of chromium. The effect of temperature is reported as the test temperature at which strength is reduced to 60% of its room-temperature value. Chromium is most effective in strengthening these chromium-molybdenum steels when it is used in amounts of 1 to 21=2%. Fig. 43 Effect of chromium content on strength. Test temperature required to reduce tensile strength and yield strength to 60% of their room-temperature values for chromium-molybdenum steels containing 0.5 to 1.0% Mo and the indicated amount of chromium

Molybdenum is an essential alloying element in ferritic steels where good creep resistance above 450 °C (840 °F) is required. Even in small amounts (0.1 to 0.5%), molybdenum increases the resistance of these steels to deformation at elevated


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temperatures. Much greater creep strength can be obtained by increasing the molybdenum level to about 1% but at the expense of greatly reduced rupture ductility (Ref 69). Additions of chromium can improve rupture ductility. Molybdenum is a carbide stabilizer and prevents graphitization. For certain ranges of stress and temperature, the dissolving of iron carbide and the concurrent precipitation of molybdenum carbide cause strain hardening in these steels. Molybdenum in amounts of 0.5% or less also minimizes temper embrittlement. Niobium and vanadium are added to improve elevated-strength properties. Vanadium is also added to some of the higher-carbon steels to provide additional resistance to tempering and to retard the growth of carbides at service temperatures. Niobium is sometimes added to these steels to increase their strength through the formation of carbides. Niobium and vanadium improve resistance to hydrogen attack, but may promote hot (reheat) cracking. Boron is added to increase hardenability. Boron can cause hot shortness and can impair toughness. Tungsten behaves like molybdenum in simple steels and has been proposed for replacing molybdenum in nuclear applications (Ref 70, 71, and 72). Thermal Exposure and Aging. Thermal exposure over time is one of the main service conditions affecting mechanical properties because the metallurgical structure of steel changes with time at temperature. For example, a ferritic matrix may be either fine or coarse-grained initially, and the carbides may vary from lamellar to completely spheroidized. With increasing time at service temperatures, the metallurgical structure slowly approaches a more stable state. For example, there may be some increase in ferrite grain size, the carbides may spheroidize, and the structure of carbon and carbon-molybdenum steels may approach the graphitized condition, with large irregular nodules of graphite in a ferrite matrix and few, if any, remaining carbides. The thermal exposure of molybdenum and molybdenum-chromium ferritic steels also contributes to complex aging phenomena, which are governed by the complicated carbide precipitation processes that occur in the steel. Figure 32 , for example, shows the sequence of carbide formations in 2.25Cr-1Mo steel. The M2C carbide (where M is primarily molybdenum) is the principal carbide for strengthening in this steel. The Mo2C first precipitates during heat treatment and/or elevated-temperature exposure. The Mo2C forms a high density of fine needles or platelets and thus contributes to strengthening by dispersion hardening. During thermal exposure, however, the unstable Mo2C carbide eventually transforms into large globular particles of M23C and η carbide. These particles are thought to have little strengthening effect, although there are some indications that M23C and η carbide present after long aging times in 21=4Cr-1Mo steel can enhance rupture strength (Ref 37). Precipitation kinetics also depend on microstructure. The strengthening carbide Mo 2C precipitates more rapidly in bainite than proeutectoid ferrite. Similarly, the Mo2C is replaced more quickly by more stable carbides in bainite than in proeutectoid ferrite. In either case, these precipitation reactions influence the strength in a similar way, regardless of whether the microstructure is bainite or proeutectoid ferrite. Spheroidization of the carbides in a steel occurs over time because spheroidized microstructures are the most stable microstructure found in steels. This spheroidization of carbides reduces strength and increases ductility. The effect of spheroidization on the rupture strength of a typical carbon-molybdenum steel containing 0.17% C and 0.42% Mo, at 480 and 540 °C (900 and 1000 °F), is shown in Fig. 44 for several initial metallurgical structures (normalized or annealed, fine or coarse grained). In these tests, the structure of the steel affected the rupture strength; for example, the stress for failure of a spheroidized structure in a given time was sometimes only half that of a normalized structure. Fig. 44 Effect of spheroidization on the rupture strength of carbon-molybdenum steel (0.17C-0.88Mn-0.20Si-0.42Mo). Source: Ref 73


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At 480 °C (900 °F), a coarse-grain normalized structure was the strongest for both short-time and long-time tests. The spheroidized structures were weaker than the normalized or annealed structures for short-time tests at both 480 and 540 °C (900 and 1000 °F). As the test time increased, the rupture values for all the structures tended to approach a common value. The rate of spheroidization depends on the initial microstructure. The slowest spheroidizing is associated with pearlitic microstructures, especially those with coarse interlamellar spacings. Spheroidizing is more rapid if the carbides are initially in the form of discrete particles, as in bainite, and even more rapid if the initial structure is martensite. Graphitization is a microstructural change that sometimes occurs in carbon or low-alloy steels subjected to moderate temperatures for long periods of time. The microstructure of carbon and carbon-molybdenum steels used for high-temperature applications such as vessels or pipes is normally composed of pearlite, which is a mixture of ferrite with some iron carbide (cementite). However, the stable form of carbon is graphite rather than cementite. Therefore, the pearlite can decompose into ferrite and randomly dispersed graphite, while the cementite will tend to disappear in these materials if they are in service long enough at metal temperatures higher than 455 °C (850 °F). This graphitization from the decomposition of pearlite into ferrite and carbon (graphite) can embrittle steel parts, especially when the graphite particles form along a continuous zone through a load-carrying member. Graphite particles that are randomly distributed throughout the microstructure cause only moderate loss of strength. Graphitization can be resisted by steels containing more than 0.7% Cr; such steels always contain at least 0.5% Mo as well, largely to impart elevated-temperature strength and resistance to temper embrittlement. Graphitization and the formation of spheroidal carbides are competing mechanisms of pearlite decomposition. The rate of decomposition is temperature dependent for both mechanisms, and the mechanisms have different activation energies. As shown in Fig. 45 , graphitization is the usual mode of pearlite decomposition at temperatures below about 550 °C (about 1025 °F), and the formation of spheroidal carbides can be expected to predominate at higher temperatures. Because graphitization involves prolonged exposure to moderate temperatures, it seldom occurs in boiling-surface tubing. Economizer tubing, steam piping, and other components that are exposed to temperatures from about 425 to 550 °C (800 to 1025 °F) for several thousand hours are more likely than boiler-surface tubing to be embrittled by graphitization. Fig. 45 Temperature-time plot of pearlite decomposition by the competing mechanisms of spheroidization and graphitization in carbon and low-alloy steels. The curve for spheroidization is for conversion of one-half of the carbon in 0.15% C steel to spheroidal carbides. The curve for graphitization is for conversion of one-half of the carbon in aluminum-deoxidized, 0.5% Mo cast steel to nodular graphite.

The heat-affected zones adjacent to welds are among the most likely locations for graphitization to occur. Figure 46 (a) shows a carbon-molybdenum steel tube that ruptured in a brittle manner along fillet welds after 13 years of service. Investigation of this failure revealed that the rupture was caused by the presence of chainlike arrays of embrittling graphite nodules (Fig. 46 b and c) along the edges of heat-affected zones associated with each of the four welds on the tube. Arrays of graphite nodules were also found in the same locations on welds in several adjacent tubes, necessitating replacement of the entire tube bank. Fig. 46 Carbon-molybdenum steel tube that ruptured in a brittle manner after 13 years of service because of graphitization at weld heat-affected zones. (a) View of tube showing dimensions, locations of welds, and rupture. (b) Macrograph showing graphitization along edges of a weld heat-affected zone (at A); this was typical of all four welds. 2×. (c) Micrograph of a specimen etched in 2% nital showing chainlike array of embrittling graphite nodules (black) at the edge of a heat-affected zone. 100×


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Decarburization is a loss of carbon from the surface of a ferrous alloy as a result of heating in a medium (for example, hydrogen) that reacts with carbon. Unless special precautions are taken, the risk of losing carbon from the surface of steel is always present in any heating to high temperatures in an oxidizing atmosphere. A marked reduction in fatigue strength is noted in steels with decarburized surfaces. The effect of decarburization is much greater on high tensile strength steels than on steels with low tensile strength. Carburization. As in the case of sulfide penetration, carburization of high-temperature alloys is thermodynamically unlikely except at very low oxygen partial pressures, because the protective oxides of chromium and aluminum are generally more likely to form than the carbides. However, carburization can occur kinetically in many carbon-containing environments. Carbon transport across continuous nonporous scales of Al 2O3 or Cr2O3 is very slow, and alloy pretreatments likely to promote such scales (for example, initially smooth surfaces or preoxidation) have generally been found to be effective in decreasing carburization attack. The suitability of carburized metal for further service can be determined by evaluating its properties and condition. The mechanical properties of the carburized layer vary markedly from those of the unaffected metal. Room-temperature ductility and toughness are decreased, and hardness is increased greatly. This deterioration is important if the carburized layer is stressed in tension because cracking is quite likely to occur. Weldability is adversely affected. Welds in carburized materials frequently show cracks because of thermal tensile stresses, even with preheating and postheating. Ductility at temperatures above 400 °C (750 °F) is usually adequate. The corrosion resistance of the low chromium-molybdenum steels commonly used for elevated-temperature applications is reduced because of the reduction in effective chromium content. For the same reason, carburized stainless steels may have a relatively low resistance to general corrosion and to intergranular corrosion, particularly while the equipment is shut down. Minor amounts of carburization do not affect creep and rupture strengths significantly. Factors Affecting Fatigue Strength. As described in the section "Creep-Fatigue Interaction" in this article, the hold times (dwell periods) and the waveform of cyclic strains influence the fatigue strength of metals at elevated temperatures. These factors affect the assessment of low-cycle fatigue and creep fatigue (Fig. 19 ). In addition, environmental effects and strain aging also influence fatigue strength. Environmental Effects. It has long been recognized that oxidation at elevated temperatures can have a marked effect (usually an acceleration) on fatigue crack initiation and growth (Ref 74, 75, 76, 77, 78, 79, and 80). In many alloys, intergranular oxidation initiates intergranular cracks at temperatures that depend on waveform or frequency but are near one-half the melting point (Ref 77). Penetration of oxygen along slip bands with subsequent localized embrittlement and cracking is another possibility. An example of the effects of environment on the fatigue strength of 21=4Cr-1Mo steel is shown in Fig. 47 . These tests were conducted in bending at a frequency of 0.05 Hz. Fig. 47 Fatigue test results of 21=4Cr-1Mo steel in sodium, air, and helium at 593 °C (1100 °F). Source: Ref 81


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Dynamic Strain Aging. In addition to the effects of environment and temperature, ferritic low-carbon and alloy steels, when subjected to inelastic deformation in certain ranges of temperature, strain, and strain rate, undergo dynamic strain aging. Dynamic strain aging, which involves the interaction of interstitials and/or carbide or nitride formers such as chromium, molybdenum, and manganese with strain-induced dislocations, has been shown to markedly influence the cyclic strain rate dependent hardening characteristics, thus affecting both the initiation and growth of fatigue cracks in ferritic materials (Ref 81). Figure 48 shows the deleterious effect of a decreasing strain rate on the fatigue strength of an annealed 21=4Cr-1Mo steel at various temperatures. Typically, strength is increased and ductility is decreased over the temperature ranges where aging occurs; therefore, both low-and high-cycle fatigue properties can be influenced accordingly. Thus, understanding fatigue, creep fatigue, environment, and strain aging interactions in the intermediate-to high-cycle life region is important. Fig. 48 Cycles to failure as a function of temperature and strain rate (continuous cycling) for various heats of isothermally annealed 21=4Cr-1Mo steel. Source: Ref 81


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Air Hardening. The amount of air hardening of the various low-alloy steels depends on composition and austenitizing temperatures (Table 5 ). Steels that harden to more than 300 HB upon air cooling from 1065 °C (1950 °F) require close control of welding operations. Gas-shielded fusion methods with nonconsumable tungsten electrodes and filler rods of parent metal are suitable for welding the 5% Cr steels having about 0.40% C. Except for the chromium-molybdenum-vanadium type, these steels are normally annealed at the mill, and no further heat treatment is necessary except stress relief after welding for the air-hardening steels. Heavy sections can be preheated. Table 5 Air-hardening characteristics of low-alloy high-temperature steels Specimens were cylinders 25 mm (1 in.) in diameter and 50 mm (2 in.) long. After air cooling from the designated temperatures they were sliced longitudinally, and the hardness determined. The values are influenced by chemical variations within the permissible limits of the specifications. Hardness, HB, after air cooling from Annealed

760 °C (1400 °F)

815 °C (1500 °F)

870 °C (1600 °F)

900 °C (1650 °F)

955 °C (1750 °F)

1010 °C (1850 °F)

1065 °C (1950 °F)

0.50Mo

137

149

149

149

163

170

187

187

1Cr-0.50Mo

137

149

149

149

170

181

187

187

1.25Cr-0.50Mo

156

149

156

179

187

229

223

212

2.25Cr-1Mo

140

149

149

235

311

311

321

285

Type 502

137

137

137

321

341

341

341

341

7Cr-0.5Mo

156

156

156

321

363

388

388

363

9Cr-0.5Mo Source: Ref 9

163

170

170

269

321

388

388

375

Steel

Elevated-Temperature Behavior of 2.25Cr-1Mo Steel The elevated-temperature behavior of 21=4Cr-1Mo steel has been studied more thoroughly than that of any other steel. The available data on annealed and normalized and tempered 21=4Cr-1Mo steel are summarized in Ref 82 and 83. The rupture strength and creep ductility of 21=4Cr-1Mo steel in various heat-treated conditions are reviewed in Ref 82. The following conclusions were reached: • The stress-rupture strength generally increases linearly with room-temperature tensile strength up to about 565 °C (1050 °F) for times up to 10,000 h • At a given strength level, tempered bainite results in higher creep strength than tempered martensite or ferrite-pearlite aggregates for temperatures up to 565 °C (1050 °F) and times up to 100,000 h. For higher temperatures and times, the ferrite-pearlite structure is the strongest • Rupture ductility generally decreases with rupture time, reaches a minimum, and then increases again. Test temperature, room-temperature tensile strength, austenitizing temperature, and impurity content increase the rate of decrease of ductility with time and cause the ductility minimum to occur at shorter times In terms of application, this steel has an excellent service record in both fossil fuel and nuclear fuel plants for generating electricity. The severe operating conditions in these plants have justified extensive studies of the behavior of 2 1=4Cr-1Mo steel under complex loading conditions and in unusual environments. This steel has become a reference against which the performance of other steels can be measured. Specifications, Steelmaking Practices, and Heat Treatments. Some of the specifications for 21=4Cr-1Mo steel in the ASME Boiler and Pressure Vessel Code are listed in Table 6 , which also includes product forms and room-temperature mechanical property requirements. For some of these specifications, composition ranges and limits differ slightly from those given in Tables 3(a) and 3(b) . Table 6 Room-temperature mechanical properties of 21=4Cr-1Mo steel in various product forms Mechanical properties Yield strength ASME specification

Grade

Product form

MPa

ksi

Ultimate tensile strength MPa

ksi

Minimu m reduction Minimum in elongation in 50 mm (2 in.), % area, %

SA-182

F22

Pipe flanges, fillings, and valves

275

40

485

70

...

...

SA-199

T22

Seamless cold-drawn tubes

170

25

415

60

30

...

SA-213

T22

Seamless ferritic alloy steel tubes

170

25

415

60

30

...


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275

40

480

01 Sep 2005

SA-217

WC9

Alloy steel castings

70

20

33

SA-333

P22

Welded and seamless pipe

205

30

415

60

30

20

SA-336

F22

Alloy steel forgings

310

45

515−755

75−110

18

25

F22a

Alloy steel forgings

205

30

415−585

60−85

20

35

SA-369

FP22

Ferritic alloy steel forged and bored pipe

205

30

415

60

20−30

...

SA-387

GR22, class 1 Chromium-molybdenum PV plate

205

30

415−585

60−85

18(a), 45

40

GR22, class 2 Chromium-molybdenum PV plate

310

45

515−690

75−100

18(a), 45

40

SA-426

CP22

Centrifugally cast ferritic alloy steel pipe

275

40

480

70

20

35

SA-542

Class 1

Chromium-molybdenum alloy steel plate

585

85

725−860

105−125

14

...

Class 2

Chromium-molybdenum alloy steel plate

690

100

790−930

115−135

15

...

(a) Elongation in 200 mm (8 in.)

In the United States, 21=4Cr-1Mo steel is normally manufactured in an electric furnace. In Japan, basic oxygen processes are used. For certain critical applications, vacuum arc remelting or electroslag remelting is appropriate. The austenitizing temperature for 21=4Cr-1Mo steel is about 900 °C (1650 °F). Heat treatments commonly employed with 21=4Cr-1Mo steel include: • Normalize and temper: Austenitize at 910 to 940 °C (1650 to 1725 °F), cool in air, temper at 580 to 720 °C (1075 to 1325 °F) • Oil quench and temper: Austenitize at 940 to 980 °C (1725 to 1800 °F), quench in oil, temper at 570 to 705 °C (1065 to 1300 °F) Short-Term Elevated-Temperature Mechanical Properties of 21=4Cr-1Mo Steel. The effects of test temperature on the tensile and yield strengths of 21=4Cr-1Mo steel are illustrated in Fig. 11 , 24 , and 49 . Data for annealed specimens and for hardened and tempered specimens are also included. The large variations in both tensile strength and yield strength with temperature and strain rate (Fig. 11 ) are caused by strain rate, temperature, and microstructure. Fig. 49 Effect of test temperatures on strength of 21=4Cr-1Mo steel. Effect of test temperature on tensile strength, yield strength, creep strength (for creep rate of 0.1 µm/m · h), and stress to rupture (for life of 100,000 h) of annealed specimens (dashed lines) and hardened and tempered specimens (solid lines) of 21=4Cr-1Mo steel. Source: Ref 2


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01 Sep 2005

The effects of elevated temperatures, elongation, and reduction in area for annealed specimens tested at standard strain rates are illustrated in Fig. 50 . Specimens tested at about 400 °C (750 °F) showed both an increase in strength and a reduction in ductility, both of which were caused by strain aging. However, the reduction in ductility was relatively small. Fig. 50 Effect of test temperature on ductility. (a) Elongation in 50 mm (2 in.) and (b) reduction in area for annealed specimens of 21=4Cr-1Mo steel tested at the indicated temperatures. Source: Ref 84


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01 Sep 2005

The effects of temperature on modulus of elasticity, shear modulus, and Poisson's ratio are shown in Fig. 12 . The modulus of elasticity diminishes from 215 GPa (31 × 106 psi) at room temperature to 140 GPa (20.3 × 106 psi) at 760 °C (1400 °F); similarly, the shear modulus diminishes from 83 GPa (12.05 × 106 psi) at room temperature to 52.4 GPa (7.6 × 106 psi) at 760 °C (1400 °F). Poisson's ratio increases from 0.288 at room temperature to 0.336 at 760 °C (1400 °F). Long Term Elevated-Temperature Mechanical Properties of 2 1=4Cr-1Mo Steel. The creep and stress-rupture behavior of annealed specimens and hardened and tempered specimens of 21=4Cr-1Mo steel are illustrated in Fig. 25 , 26 , 27 , and 49 . With regard to rupture life and creep rate, the hardened and tempered specimens were able to withstand higher stresses than the annealed specimens. The ductility exhibited by stress-rupture specimens can be roughly correlated with stress level or rupture life. In general, specimens tested at high stress levels have short rupture lives, and such specimens exhibit greater reduction in area than similar specimens tested at lower stress levels. These data show considerable scatter but no evidence of brittle behavior by this steel. The relaxation behavior of 21=4Cr-1Mo steel is illustrated in Fig. 51 . Fig. 51 Relaxation behavior of 21=4Cr-1Mo steel. Specimens were stressed to level indicated on ordinate of graph and exposed to elevated temperature for indicated duration; remaining stress indicated on graph. Source: Ref 84

Long-term exposure to elevated temperature can reduce the room-temperature and elevated-temperature properties of 21=4Cr-1Mo steel. Some of these effects are illustrated in Fig. 6 (a), 7 (a), 52 , and 53 . Figure 7 (a) shows the changes in room-temperature tensile properties caused by exposure (without stress) to elevated temperatures. Fig. 52 Effect of exposure to elevated temperature on the strength of 21=4Cr-1Mo steel. Variation in tensile and yield strengths of two different heats of 21=4Cr-1Mo steel after exposure (without stress) to test temperature of 455 °C (850 °F)


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01 Sep 2005

Fig. 53 Effect of exposure to elevated temperature on the stress-rupture behavior of 21=4Cr-1Mo steel. Variation in rupture life for specimens of annealed 21=4Cr-1Mo steel exposed to various elevated temperatures for the durations indicated. After aging, all specimens were stressed to 140 MPa (20 ksi) and tested at 565 °C (1050 °F).

Figure 52 shows the effect of variations in aging time (without stress) at 455 °C (850 °F) on the ultimate tensile and yield strengths of two heats of 21=4Cr-1Mo steel tested at the same temperature. The difference in strength between these two heats was observed even before the tests; the differences were probably caused by variations in composition and microstructure. The same factors account for strength changes during aging because they affect both the size and distribution of carbides in the steel. As shown in Fig. 53 , prolonged aging without stress at 565 °C (1050 °F) can reduce time to rupture for annealed 21=4Cr-1Mo steel. Similarly, the data in Fig. 6 (a) show that prolonged exposure to high temperatures without stress substantially reduces stress to rupture in a fixed time. The amount of reduction in stress to rupture is greatest for exposure at 480 °C (900 °F). Elevated-Temperature Fatigue Behavior of 21=4Cr-1Mo Steel. The results of strain-controlled fatigue tests at 425, 540, and 595 °C (800, 1000, and 1100 °F) on specimens of annealed 21=4Cr-1Mo steel are shown in Fig. 54 . Within this range, the test temperature had relatively little effect on the number of cycles to failure. Other strain-controlled fatigue tests (Fig. 48 ) have shown that reducing the carbon content to 0.03% decreases the fatigue strength. Furthermore, because of variations in strain-aging effect, specimens from one heat with a higher carbon content ran longer at 425 °C (800 °F) than at 315 °C (600 °F). Fig. 54 Effect of elevated temperature on strain-controlled fatigue behavior of annealed 2 1=4Cr-1Mo steel. Strain rate was greater than 4 mm/m · s. Source: Ref 84


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01 Sep 2005

The crack growth rate data shown in Fig. 55 and 56 were obtained from precracked specimens subjected to cyclic loading at a constant maximum load. Crack extension was measured at intervals during testing. The stress intensity factor range increased as crack length was increased. Figure 55 illustrates the increase in crack growth rate with increasing test temperature. The data in Fig. 56 indicate that in elevated-temperature tests at a given stress intensity factor range, crack growth rate increases as cyclic frequency is decreased. These fracture mechanics data can be applied to the design of structural components that may contain undetected discontinuities or that may develop cracks in service. Fig. 55 Effect of temperature on fatigue crack growth rate. Variations in fatigue crack growth rate with test temperature for specimens of 21=4Cr-1Mo steel tested in air. Stress ratio was 0.05; cyclic frequency was 400 per minute. Source: Ref 85

Fig. 56 Effect of cyclic frequency on fatigue crack growth rate. Variations in fatigue crack growth rate with cyclic frequency for specimens of 21=4Cr-1Mo steel tested in air. Stress ratio was 0.05. (a)Tested at 510 °C (950 °F). (b)Tested at 595 °C (1100. °F). Source: Ref 85


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01 Sep 2005

The introduction of a holding period at the peak strain of each fatigue cycle reduces fatigue life as described in the section "Creep-Fatigue Interaction." From studies conducted on annealed 2.25Cr-1Mo steel in air (Ref 6, 81, 86), it is possible to conclude the following (Ref 30): • Compressive hold periods are more damaging than tensile hold periods, and hold periods imposed on the tension-going side of the hysteresis loop are more damaging in terms of reduced cycle life than hold periods on the compression-going side • Linear damage summation of fatigue and creep damage does not sum to a unique value The fact that the damage sums are less than 1 indicates apparent creep-fatigue interaction, but because the values shown do not trend toward a unique value, the linear damage summation method is highly questionable for data extrapolation. The primary reason that the damage sums are less than 1 is that significant environmental interaction or corrosion fatigue occurs in air. This oxidation can substantially reduce the time for crack initiation in smooth bar tests, depending on waveform, and is not adequately accounted for by the simple linear damage summation of fatigue and creep damage fractions. Environmental interaction is discussed in the section "Factors Affecting Fatigue Strength" in this article. Properties of Welds in 21=4Cr-1Mo Steel. Welding is often required in the fabrication of pressure vessels, boilers, heat exchangers, and similar structures for use at elevated temperatures in power plants, refineries, chemical-processing plants, and similar applications. Therefore, in evaluating materials for these structures, it is important to consider the mechanical properties of welded joints. In one investigation, the elevated-temperature tensile and creep-rupture properties of weldments in 21=4Cr-1Mo steel were measured (Ref 87, 88). Specimens were cut from the weld metal and the base metal; other specimens had transverse welds. All specimens were tempered at 705 °C (1300 °F) before testing. In all these tests, the weld metal was stronger than either the base metal or the specimens containing transverse welds. Specimens with transverse welds invariably fractured in the base metal. The high strength of the weld metal relative to that of the base metal was attributed of differences in microstructure. The base metal, which had been normalized and tempered, contained more ferrite and less bainite than the weld metal. In these tests, the base metal was the weakest part of the welded structure.

Thermal Expansion and Conductivity Because of their higher thermal conductivity and lower thermal coefficient of expansion, ferritic steels may be more desirable than austenitic steels when thermal cycling occurs in service. Figures 57 and 58 indicate the thermal conductivity and expansion coefficient for carbon and low-alloy steels as a function of temperature. Fig. 57 Thermal conductivity of carbon and low-alloy steels at various temperatures


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01 Sep 2005

Fig. 58 Coefficients of thermal expansion for carbon and low-alloy steels at various temperatures. These are not mean values of the coefficient over a range of temperatures.


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01 Sep 2005

ACKNOWLEDGMENT ASM INTERNATIONAL would like to thank R.L. Klueh of Oak Ridge National Laboratory for providing literature to update this article. Thanks are also extended to Joseph Conway of Mar-Test, Inc., and R.W. Swindeman and C.R. Brinkman of Oak Ridge National Laboratory for their review. REFERENCES 1. G.V. Smith, Evaluation of the Elevated Temperature Tensile and Creep-Rupture Properties of 1=2Cr-1=2Mo, 1Cr-1=2Mo, and 11=4Cr-1=2Mo-Si Steels, DS 50, American Society for Testing and Materials, 1973 2. G.V. Smith, Supplemental Report on the Elevated-Temperature Properties of Chromium-Molybdenum Steels (An Evaluation of 21=4Cr-1Mo Steel), DS 6 S2, American Society for Testing and Materials, March 1971 3. G.S. Sangdahl and H.R. Voorhees, Quenched-and-Tempered 21=4Cr-1Mo Steel at Elevated Temperatures⎯Tests and Evaluation, in 21=4Chrome-1 Molybdenum Steel in Pressure Vessels and Piping, American Society of Mechanical Engineers, 1972 4. G.S. Sangdahl and M. Semchyshen, Ed., Application of 21=4Cr-1Mo for Thick-Wall Pressure Vessels, STP 755, American Society for Testing and Materials, 1982 5. Low Carbon and Stabilized 2 1=4% Chromium 1% Molybdenum Steels, American Society for Metals, 1973 6. C.R. Brinkman et al., Time-Dependent Strain-Controlled Fatigue Behavior of Annealed 21=4Cr-1Mo Steel for Use in Nuclear Steam Generator Design, J. Nucl. Mater., Vol 62, 1976, p 181−204 7. V.K. Sikka, "Development of a Modified 9Cr-1Mo Steel for Elevated Temperature Service," in Proceedings of Topical Conference on Ferritic Alloys for Use in Nuclear Energy Technologies, The Metallurgical Society of AIME, 1984, p 317−327 8. R. Viswanathan, Strength and Ductility of CrMoV Steels in Creep at Elevated Temperatures, ASTM J. Test. and Eval., Vol 3 (No. 2), 1975, p 93−106


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9. R. Viswanathan and R.I. Jaffee, Toughness of Cr-Mo-V Steels for Steam Turbine Rotors, ASME J. Eng. Mater. Tech., Vol 105, Oct 1983, p 286−294 10. R. Crombie, High Integrity Ferrous Castings for Steam Turbines⎯Aspects of Steel Development and Manufacture, Mater. Sci. Tech., Vol 1, Nov 1985, p 986−993 11. J.A. Todd et al., New Low Chromium Ferritic Pressure Vessel Steels, in Mi-Con 86: Optimization of Processing, Properties, and Service Performance Through Microstructural Control, STP 979, American Society for Testing and Materials, 1986, p 83−115 12. R.G. Baker and J. Nutting, J. Iron Steel Inst., Vol 192, 1959, p 257−268 13. T. Ishiguro et al., Research on Chrome Moly Steels, R.A. Swift, Ed., MPC-21, American Society of Mechanical Engineers, 1984, p 43−51 14. V.K. Sikka, M.G. Cowgill, and B.W. Roberts, Creep Properties of Modified 9Cr-1Mo Steel, in Conference on Ferritic Alloys for Use in Nuclear Energy Technologies, American Institute of Mining, Metallurgical and Petroleum Engineers, 1984, p 413−423 15. V.K. Sikka, G.T. Ward, and K.C. Thomas, in Ferritic Steels for High Temperature Applications, American Society for Metals, 1982, p 65−84 16. R.L. Klueh and R.W. Swindeman, The Microstructure and Mechanical Properties of a Modified 2.25Cr-1Mo Steel, Metall. Trans. A, Vol 17A, 1986, p 1027−1034 17. R.L. Klueh and A.M. Nasreldin, Metall. Trans. A, Vol 18A, 1987, p 1279−1290 18. W.B. Jones, Effects of Mechanical Cycling on the Substructure of Modified 9Cr-1Mo Ferritic Steel, in Ferritic Steels for High-Temperature Applications, A.K. Khare, Ed., American Society for Metals, 1983, p 221−235 19. J.L. Handrock and D.L. Marriot, Cyclic Softening Effects on Creep Resistance of Bainitic Low Alloy Steel Plain and Notched Bars, in Properties of High Strength Steels for High-Pressure Containments, E.G. Nisbett, Ed., MPC-27, American Society of Mechanical Engineers, 1986 20. R.W. Swindeman, Cyclic Stress-Strain-Time Response of a 9Cr-1Mo-V-Nb Pressure Vessel Steel at High Temperature, in Low Cycle Fatigue, STP 942, American Society for Testing and Materials, 1987, p 107−122 21. S. Kim and J.R. Weertman, Investigation of Microstructural Changes in a Ferritic Steel Caused by High Temperature Fatigue, Metall. Trans. A, Vol 19A, 1988, p 999−1007 22. R.L. Klueh and R.E. Oakes, Jr., High Strain Rate Tensile Properties of Annealed 21=4Cr-1Mo Steel, J. Eng. Mater. Technol., Vol 98, Oct 1976, p 361−367 23. Digest of Steels for High Temperature Service, 6th ed., The Timken Roller Bearing Company, 1957 24. M. Prager, Factors Influencing the Time-Dependent Properties of Carbon Steels for Elevated Temperature Pressure Vessels, MPC 19, American Society of Mechanical Engineers, 1983, p 12, 13 25. R.M Goldhoff, Stress Concentration and Size Effects in a CrMoV Steel at Elevated Temperatures, Joint International Conference on Creep, Institute of Mechanical Engineers, London, 1963 26. R. Viswanathan and C.G. Beck, Effect of Aluminum on the Stress Rupture Properties of CrMoV Steels, Met. Trans. A, Vol 6A, Nov 1975, p 1997−2003 27. "Aerospace Structural Metals Handbook," AFML-TR-68-115, Army Materials and Mechanics Research Center, 1977 28. J.W. Freeman and H. Voorhees, in Relaxation Properties of Steels and Superstrength Alloys at Elevated Temperatures, STP 187, American Society for Testing and Materials 29. H.R. Voorhees and M.J. Manjoine, Compilation of Stress-Relaxation Data for Engineering Alloys, DS 60, American Society for Testing and Materials, 1982 30. C.R. Brinkman, High-Temperature Time-Dependent Fatigue Behavior of Several Engineering Structural Alloys, Int. Met. Rev., Vol 30 (No. 5), 1985, p 235−258 31. D.A. Miller, R.H. Priest, and E.G. Ellison, A Review of Material Response and Life Prediction Techniques Under Fatigue-Creep Loading Conditions, High Temp. Mater. Proc., Vol 6 (No. 3 and 4), 1984, p 115−194 32. R. Viswanathan, Damage Mechanisms and Life Assessment of High-Temperature Components, ASM INTERNATIONAL, 1989 33. Thermal Fatigue of Materials and Components, STP 612, American Society for Testing and Materials, 1976 34. M.A. Grossman and E.C. Bain, Principles of Heat Treatment, 5th ed., American Society for Metals, 1964 35. The Generation of Isochronous Stress-Strain Curves, A.O. Schaefer, Ed., American Society of Mechanical Engineers, 1972 36. G.V. Smith, Evaluation of the Elevated Temperature Tensile and Creep-Rupture Properties of 3−9% Chromium-Molybdenum Steels, DS 58, American Society for Testing and Materials, 1975 37. R.L. Klueh, Interaction Solid Solution Hardening in 2.25Cr-1Mo Steel, Mater. Sci. Eng., Vol 35, 1978, p 239−253 38. R. Viswanathan and R.D. Fardo, Parametric Techniques for Extrapolating Rupture Ductility, in Ductility and Toughness Considerations in Elevated Temperature Service, G.V. Smith, Ed., MPC-8, American Society of Mechanical Engineers,


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1978 39. J. Gutzeit, High Temperature Sulfidic Corrosion of Steels, in Process Industries Corrosion⎯The Theory and Practice, National Association of Corrosion Engineers, 1986 40. H.F. McConomy, High-Temperature Sulfidic Corrosion in Hydrogen-Free Environment, Proc. API, Vol 43 (III), 1963, p 78−96 41. E.B. Backensto, R.D. Drew, and C.C. Stapleford, High Temperature Hydrogen Sulfide Corrosion, Corrosion, Vol 12 (No. 1), 1956, p 6t−16t 42. G. Sorell and W.B. Hoyt, Collection and Correlation of High Temperature Hydrogen Sulfide Corrosion Data, Corrosion, Vol 12 (No. 5), 1956, p 213t−234t 43. C. Phillips, Jr., High Temperature Sulfide Corrosion in Catalytic Reforming of Light Naphthas, Corrosion, Vol 13 (No. 1), 1957, p 37t−42t 44. G. Sorell, Compilation and Correlation of High Temperature Catalytic Reformer Corrosion Data, Corrosion, Vol 14 (No. 1), 1958, p 15t−26t 45. W.H. Sharp and E.W. Haycock, Sulfide Scaling Under Hydrorefining Conditions, Proc. API, Vol 39 (III), 1959, p 74−91 46. J.D. McCoy and F.B. Hamel, New Corrosion Data for Hydrosulfurizing Units, Hydrocarbon Process., Vol 49 (No. 6), 1970, p 116−120 47. J.D. McCoy and F.B. Hamel, Effect of Hydrosulfurizing Process Variables on Corrosion Rates, Mater. Prot. Perform., Vol 10 (No. 4), 1971, p 17−22 48. A.S. Couper and J.W. Gorman, Computer Correlations to Estimate High Temperature H2S Corrosion in Refinery Streams, Mater. Prot. Perform., Vol 10 (No. 1), 1971, p 31−37 49. D. Warren, Hydrogen Effects on Steel, in Process Industries Corrosion, National Association of Corrosion Engineers, 1986, p 21−30 50. G.R. Odette, Conference Proceedings on Materials for Coal Conversion and Utilization, National Bureau of Standards, 1982 51. R.O. Ritchie et al., J. Mater. Energy Sys., Vol 6 (No. 3), p 151−162 52. G.A. Nelson, Metals for High Pressure Hydrogenation Plants, Trans. ASME, Vol 73, 1951, p 205−213 53. I. Masaoka et al., Hydrogen Attack Limit of 21=4Cr-1Mo Steel, in Current Solutions to Hydrogen Problems in Steel, American Society for Metals, 1982, p 247 54. P.G. Shewmon et al., On the Nelson Curve for 21=4Cr-1Mo Steel, in Research on Chrome-Moly Steels, MPC-21, American Society of Mechanical Engineers, 1984, p 1−8 55. R.L Schuyler III, Hydrogen Blistering of Steel in Anhydrous Hydrofluoric Acid, Mater. Perform., Vol 18 (No. 8), 1979, p 9−16 56. G. Herbsleb et al., Occurrence and Prevention of Hydrogen Induced Stepwise Cracking and Stress Corrosion Cracking of Low Alloy Pipeline Steels, Corrosion, Vol 37 (No. 5), 1981, p 247−255 57. J.D. Baird et al., Strengthening Mechanisms in Ferritic Creep Resistant Steels, in Creep Strength in Steel and High Temperature Alloys, The Metals Society, in 1974, p 207−216 58. J.D. Baird and A. Jamieson, J. Iron Steel Inst., Vol 210, 1972, p 841 59. J.D. Baird and A. Jamieson, J. Iron Steel Inst., Vol 210, 1972, p 847 60. B.B. Argent et al., J. Iron Steel Inst., Vol 208, 1970, p 830−843 61. G. Krauss, Principles of Heat Treatment of Steel, American Society for Metals, 1980 62. F.B. Pickering, Physical Metallurgy and the Design of Steels, Applied Science, 1978 63. R. Viswanathan, Strength and Ductility of 21=4Cr-1Mo Steels in Creep, Met. Tech., June 1974, p 284−293 64. J. Orr, F.R. Beckitt, and G.D. Fawkes, The Physical Metallurgy of Chromium-Molybdenum Steels for Fast Reactor Boilers, in Ferritic Steels for Fast Reactor Steam Generators, S.F. Pugh and E.A. Little, Ed., British Nuclear Energy Society, 1978, p 91 65. R.L Klueh, J Nucl. Mater., Vol 68, 1977, p 294 66. J. Ewald, et al., Over 30 Years Joint Long-Term Research on Creep Resistant Materials in Germany, in Advances in Material Technology for Fossil Power Plants, R. Viswanathan and R.I. Jaffee, Ed., ASM INTERNATIONAL, 1987, p. 33−39 67. A. Krisch, Jernkontorets Ann., Vol 155, 1971, p 323-331 68. G.V. Smith, Properties of Metals at Elevated Temperatures, McGraw-Hill, 1950, p 231 69. J.D. Baird, Jernkontorets Ann., Vol 151, 1971, p 311−321 70. R.L. Klueh and P.J. Maziasz, Reduced-Activation Ferritic Steels: A Comparison With Cr-Mo Steels, J. Nucl. Mater., Vol 155−157, 1988, p 602−607 71. R.L. Klueh and E.E. Bloom, The Development of Ferritic Steels for Fast Induced-Radioactive Decay for Fusion Reactor


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Applications, in Nuclear Engineering and Design/Fusion 2, North-Holland, 1985, p 383−389 72. R.L. Klueh and P.J. Maziasz, Low-Chromium Reduced-Activation Ferritic Steels, in Reduced-Activation Materials for Fusion Reactors, STP 1046, American Society for Testing and Materials, to be published 73. S.H. Weaver, The Effect of Carbide Spheroidization Upon the Rupture Strength and Ductility of Carbon Molybdenum Steel, Proc. ASTM, Vol 46, 1946, p 856−866 74. L.F. Coffin, Metall. Trans., Vol 3, 1972, p 1777−1788 75. L.A. James, J. Eng. Mater. Technol., Vol 98, July 1976, p 235−243 76. M. Gell and G.R. Leverant, in Fatigue at Elevated Temperatures, STP 520 American Society for Testing and Materials, 1973, p 37−66 77. J.C. Runkle and R.M. Pelloux, in Fatigue Mechanisms, STP 675, J.T. Fong, Ed., American Society for Testing and Materials, 1979, p 501−527 78. D.J. Duquette, Environmental Effects I: General Fatigue Resistance and Crack Nucleation in Metals and Alloys, in Fatigue and Microstructure, American Society for Metals, 1979, p 335−363 79. H.L. Marcus, Environmental Effects II: Fatigue-Crack Growth in Metals and Alloys, in Fatigue and Microstructure, American Society for Metals, 1979, p 365−383 80. P. Marshall, in Fatigue and High Temperature, R.P. Skelton, Ed., Applied Science, 1983, p 259−303 81. C.R. Brinkman et al., Time-Dependent Strain-Controlled Fatigue Behavior of Annealed 21=4Cr-Mo Steel for Use in Nuclear Steam Generator Design,J. Nucl. Mater., Vol 62, 1976, p 181−204 82. R. Viswanathan, Strength and Ductility of 21=4Cr-1Mo Steels in Creep at Elevated Temperatures, Met. Technol., June 1974, p 284−294 83. G.V. Smith, Elevated Temperature Strength and Ductility of Q&T 21=4Cr-1Mo Steel, in Current Evaluation of 21=4Cr-1Mo Steel in Pressure Vessels and Piping, American Society of Mechanical Engineers, 1972 84. M.K. Booker, T.L. Hebble, D.O. Hobson, and C.R. Brinkman, Mechanical Property Correlations for 21=4Cr-1Mo Steel in Support of Nuclear Reactor Systems Design, Int. J. Pressure Vessels Piping, Vol 5, 1977 85. C.R. Brinkman, W.R. Corwin, M.K. Booker, T.L. Hebble, and R.L. Klueh, "Time Dependent Mechanical Properties of 21=4Cr-1Mo Steel for Use in Steam Generator Design," ORNL-5125, Oak Ridge National Laboratory, 1976 86. J.J. Burke and V. Weiss, in Fatigue Environment and Temperature Effects, Plenum Press, 1983, p 241−261 87. R.L. Klueh and D.A. Canonico, Microstructure and Tensile Properties of 21=4Cr-1Mo Steel Weldments With Varying Carbon Contents, Weld. J. (Research Supplement), Sept 1976 88. R.L. Klueh and D.A. Canonico, Creep-Rupture Properties of 21=4Cr-1Mo Steel Weldments With Varying Carbon Content, Weld. J. (Research Supplement), Dec 1976


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Effect of Neutron Irradiation on Propert...

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Service Characteristics of Carbon and Low-Alloy Steels Effect of Neutron Irradiation on Properties of Steels*

[*Research sponsored by the Office of Fusion Energy, U.S. Department of Energy, under contract DE-AC05-840R21400 with the Martin Marietta Energy Systems, Inc.]

R.L. Klueh, Metals and Ceramics Division, Oak Ridge National Laboratory DAMAGE TO STEELS from neutron irradiation affects the properties of steels and is an important factor in the design of safe and economical components for fission and fusion reactors. Damage occurs when high-energy neutrons displace metal atoms from their normal lattice positions to form interstitials and vacancies. It is the disposition of these defects that influences properties during and after irradiation. In addition to the formation of vacancies and interstitials, transmutation reactions can also occur when neutrons are absorbed by the atoms of an irradiated alloy. These transmutation reactions produce new metal atoms and gas atoms of hydrogen and/or helium within the alloy matrix. Of these various transmutation by-products, transmutation helium is considered the most significant in exacerbating property changes. The effects of damage caused by neutron irradiation include swelling (volume increase), irradiation hardening, and irradiation embrittlement (the influence of irradiation hardening on fracture toughness). These effects are primarily associated with high-energy (>0.1 MeV) neutrons. Consequently, irradiation damage from neutrons is of considerable importance in fast reactors, which produce a significant flux of high-energy neutrons during operation. However, irradiation damage from neutrons is also a factor in commercial light-water reactors, even though neutrons in a light-water reactor are moderated to reduce their energy (most neutrons in the spectrum of these reactors are thermal neutrons with energies much less than 1 eV). Such reactors produce a small flux of high-energy neutrons, and until recently, these neutrons were the only ones considered to cause the irradiation effects observed in power reactors. However, as discussed in the section "Irradiation Embrittlement" in this article, recent observations have indicated that thermal neutrons can also cause irradiation effects. Therefore, material damage from neutron irradiation is important not only in fast reactors, such as the experimental Fast Flux Test Facility (FFTF) in Hanford, Washington, and the Super Phenix fast-breeder electric power reactor in France, but also in the many commercial light-water power reactors. In addition, the future of economically viable fusion reactors may also depend on the development of irradiation-resistant alloys. This article discusses the effects of high-energy neutrons on steels, with particular emphasis on the steels listed in Table 1 . For the pressure vessels of light-water reactors the manganese-molybdenum-nickel ferritic steels (ASTM A 302-B and A 533-B) are commonly used. These steels are quenched and tempered, which produces a tempered martensite and/or tempered bainite microstructure. Austenitic steels such as type 316 stainless steel are proposed for fusion reactors and are used in fast reactors for fuel cladding, ducts, and other structural components. These steels are used in either the solution-annealed or the 20% cold-worked condition. Special irradiation-resistant austenitic steels have been developed for these applications. An example of such a new steel is the prime candidate alloy (PCA) for fusion (Table 1 ). For both the fast-breeder reactor and fusion reactors, chromium-molybdenum ferritic are being considered. Of special interest are the 9Cr-1MoVNb and 12Cr-1MoVW steels; 21=4Cr-1Mo steel is also considered for fusion reactors. These steels are used in a normalized and tempered condition, which gives a tempered martensite microstructure in the 9Cr-1MoVNb and 12Cr-1MoVW steels and a tempered bainite microstructure in the 21=4Cr-1Mo steel. Table 1 Typical compositions for steels of interest for nuclear reactor applications Chemical composition, wt% Steel

Cr

Ni

Mo

Mn

Si

C

V

Nb

W

Ti

N

316 stainless steel

18.0

13.0

2.5

2.0

0.8

0.05

...

...

...

0.05

0.05

PCA steel

14.0

16.0

2.5

1.7

0.4

0.05

...

...

...

0.25

...

A302

0.1

0.2

0.5

1.3

...

0.2

...

...

...

...

...

A533

...

0.5

0.5

1.3

...

0.2

...

...

...

...

...

2.25

0.25

1.0

0.5

0.3

0.12

...

...

...

...

...

Austenitic stainless steels

Ferritic steels

1

2 =4Cr-1Mo


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01 Sep 2005

9Cr-1MoVNb

9.0

0.2

1.0

0.5

0.3

0.10

0.2

0.08

0.5

...

0.05

12Cr-1MoVW

12.0

0.5

1.0

0.5

0.4

0.20

0.3

...

...

...

...

Irradiation Damage Processes The current understanding of neutron irradiation effects has been obtained from studies on materials irradiated in fission test reactors, such as the FFTF, the Experimental Breeder Reactor (EBR-II), and the High Flux Isotope Reactor (HFIR) (Ref 1). Irradiation is described in terms of the flux of neutrons striking the steel being irradiated, which is measured as the number of neutrons per square meter per second (n/m2 · s), and the fluence, which is the time-integrated flux in neutrons per square meter (n/m2). A typical flux for a fast reactor is ~5 × 1019 n/m2 · s. Displacement Damage. When a steel is irradiated in a high-energy neutron field, neutrons collide with atoms in the material and displace them from their lattice positions (Ref 2, 3). The first atom struck and displaced by a neutron is termed a primary knock-on atom. When this primary knock-on atom recoils from the impact, it collides with other atoms, which in turn recoil and collide with still other atoms. Therefore, an incoming neutron can produce a displacement cascade, by which a large number of atoms are displaced from lattice sites. The displaced atoms of a displacement cascade move into interstitial positions (termed interstitials) and leave behind a like number of vacant sites (vacancies). This displacement of atoms by irradiation is described in terms of displacements per atom (dpa), which is a measure of the average number of times an atom is displaced from its lattice position. The dpa can be calculated from the neutron fluence received by the steel (Ref 2, 3). During 1 year of fast reactor operation, each atom in stainless steel is typically displaced more than 30 times (30 dpa/yr). For a light-water reactor, the displacement rate in steel is about 0.03 dpa/yr. For a fusion reactor, displacement rates of up to 60 dpa/yr might be expected. The disposition of the defects⎯interstitials and vacancies⎯determines the effect of the atomic displacements on the properties of an irradiated material. Although the average number of atomic displacements is described by the dpa unit, only a fraction of these displacements produce damage and property effects. In general, most (typically 95 to 99%) of the displaced atoms from a displacement cascade recombine with a vacancy. This is because the interstitials and vacancies produced in a displacement cascade are near each other and have a strong likelihood of recombining. Therefore, displacement damage from neutron irradiation occurs from only a small portion of the atomic displacements. The interstitials and vacancies from this portion of the displacements do not recombine but instead migrate to sinks, where they are absorbed or accumulate. Sinks include surfaces, grain boundaries, dislocations, and existing cavities. This migration of defects can also result in the formation of defect clusters; those consisting of interstitials can evolve into dislocation loops, while vacancy clusters can develop into microvoids or cavities (Fig. 1 ). Solute clusters can also form under certain conditions. Fig. 1 Cavities (indicated as the white rectangles and circles) formed in type 316 stainless steel irradiated to 60 dpa at 600 °C (1110 °F) in the HFIR. Courtesy of P.J. Maziasz, Oak Ridge National Laboratory


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The type of cluster defect that forms depends on the irradiation temperature. Below about 0.35 Tm (where Tm is the melting point of the irradiated material in degrees kelvin), interstitials are mobile relative to vacancies, and the interstitials combine to form dislocation loops. This gives rise to an increase in strength and a decrease in ductility. Vacancies become increasingly mobile above 0.35 Tm, and a dislocation and cavity structure results (Fig. 1 ). This microstructure occurs because certain sinks have a bias and do not accept vacancies and interstitials equally (Ref 2, 3). If all sinks accepted both defects equally, the vacancies and interstitials would annihilate at a sink, and no swelling would result. However, within a grain, interstitials are accepted preferentially by dislocations. This leaves an excess of vacancies to be absorbed by cavities, giving rise to the observed swelling. Finally, at high irradiation temperatures (greater than about 0.6 Tm), defect clusters are unstable. That is, the high equilibrium vacancy concentration and rapid diffusion lead to vacancy-interstitial recombination, which thus reduces the number of defects and the effects of displacement damage on properties. However, at temperatures ¾0.5 Tm, any transmutation helium produced during irradiation can lead to problems. Cavities. Two types of cavities form during irradiation⎯bubbles and voids. Bubbles contain gas at a pressure greater than or equal to the surface tension pressure. Voids have internal gas pressure below the equilibrium pressure. The origin of gas in irradiated material is described in the section "Transmutation Helium" in this article. Radiation-Induced Segregation. Because certain alloying elements can be preferentially associated with or rejected by vacancies or interstitials, such elements can be transported to or from sinks when the defects migrate. This radiation-induced segregation can cause detrimental effects and must be considered when developing alloys for irradiation resistance (Ref 2, 3). Transmutation Helium. In addition to displacement damage, a neutron can be absorbed by an atom of the irradiated alloy, resulting in a transmutation reaction that produces a new metal atom and hydrogen and/or helium gas atoms within the alloy being irradiated. Indications are that small amounts of new metal atoms have little effect on properties. Hydrogen will have little effect on properties because, at the operating temperatures of most reactors (250 to 550 °C, or 480 to 1020 °F), it should readily diffuse from the alloy. However, any helium produced can affect the properties; it is relatively insoluble in metals and will therefore be incorporated into the bubbles or voids that can form within the matrix and on grain boundaries and precipitate interfaces. The displacement damage formed in fusion and fission reactors is similar, and fusion damage can be simulated by fission reactor irradiations. However, the much higher energy of the neutrons (up to 14 MeV) produced in a fusion reactor will lead to much more transmutation helium than occurs in most fission irradiations. Such a simultaneous development of displacement damage and helium can affect both the swelling behavior and the mechanical properties. Much recent research has been directed at determining the effect of helium on the properties of irradiated candidate structural alloys and on developing alloys that will withstand these effects (Ref 3).

Void Swelling


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As mentioned earlier in the section "Displacement Damage" in this article, the bias of dislocations for interstitials causes an excess of vacancies to agglomerate at cavities (Fig. 1 ), which thus causes a volume increase or swelling of the irradiated alloy (bias−driven void swelling). Void swelling (measured as ∆V/V, where ∆V is the change in volume of the irradiated material and V is the original volume) of several tens of percent is observed in some stainless steels. Large amounts of swelling cannot be tolerated in a reactor component, and considerable effort has been directed toward the development of swelling-resistant alloys for use in fast-breeder and fusion reactors. Void swelling is unimportant for light-water power reactors because of the low flux of high-energy neutrons in the neutron spectrum of such reactors. Swelling in Austenitic Stainless Steels. Irradiations to fluences that produce displacement damage greater than 100 dpa have been conducted on types 304 and 316 stainless steel, and the swelling can be described as a function of temperature and fluence (Ref 2, 3), as illustrated in Fig. 2 . For constant fluence, a peak swelling temperature is observed. At a constant temperature, there is an incubation time, after which swelling develops slowly with a power-law dependence on fluence. This transient regime is eventually replaced by a rapid-swelling regime characterized by a linear dependence (steady-state swelling) on fluence. Swelling in the steady-state regime for austenitic steels occurs at approximately 1%/dpa; the steady-state rate is essentially independent of composition and fabrication variables and is a weak function of temperature, irradiation rate, and stress (Ref 4, 5). Fig. 2 Effect of temperature and neutron fluence on the swelling behavior of type 316 stainless steel irradiated in a fast reactor (EBR-II). Source: Ref 3

The high rate of swelling in the steady-state regime (~1%/dpa) suggested that the structural lifetime could be increased only by extending the transient regime (Ref 4, 5). Voids nucleate during the transient period, while their growth and coalescence occur during steady state. Void nucleation is aided by small amounts of dissolved gases (oxygen and nitrogen) or gases formed by transmutation reactions (helium). Gas can combine with irradiation-induced vacancies to form bubbles. These bubbles collect vacancies until a critical radius for a void is reached, after which growth is bias driven. If the conversion from bubbles to voids can be inhibited, the transient stage can be extended, and swelling resistance is improved.


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Although void swelling occurs because of the slight bias of dislocations as sinks for interstitials, a very high dislocation density can provide sufficient sinks for both vacancies and interstitials, where they can recombine and annihilate. This information has led to the use of 20%-cold-worked type 316 stainless steel for fast-breeder reactor fuel cladding. Small amounts of titanium were found to extend the transient regime (Ref 6). Titanium was thought to getter oxygen and other dissolved gases and thus reduce the nucleation rate, which was delayed until small amounts of transmutation helium could aid nucleation. The addition of titanium and niobium to form carbides and phosphorus to form phosphides has a similar effect on transient swelling (Ref 4). In this case, the fine dispersions of stable precipitate particles trap helium at their interfaces, forming a high density of small bubbles of these locations. Because the irradiation-produced vacancies are collected by a larger number of cavities, the formation of the critical radius for void growth is delayed, and the transient regime is extended. The high density of cavities can also become the dominant sink for both vacancies and interstitials and therefore sites for recombination. Such alloying techniques have been applied in the breeder reactor and fusion reactor alloy development programs in an attempt to develop irradiation-resistant stainless steels, such as the PCA steel for fusion (Table 1 ). In a fusion reactor, the large amounts of helium generated simultaneously with displacement damage will give rise to irradiation effects different from those observed in a breeder reactor. The information available for large helium concentrations at relevant damage rates is from comparisons of the behavior of type 316 stainless steel irradiated in fast reactors, such as the EBR-II, and in a mixed-spectrum reactor, such as the HFIR. In a mixed-spectrum reactor, the neutron spectrum contains both thermal and fast neutrons. When a nickel-containing steel is irradiated in the HFIR, the thermal neutrons react with 58Ni to form helium, while the fast neutrons produce displacement damage just as they do in a fast reactor. Natural nickel contains approximately 68% 58Ni. Recent evidence indicates that increasing the He:dpa ratios from approximately 0.5 (EBR-II) to approximately 60 (HFIR) (the He concentration is in atomic parts per million, appm) may shorten or extend the transition regime, depending on the heat chemistry and the thermomechanical treatment (Ref 4). An extended transition regime is associated with a high density of bubbles, the inhibition of radiation-induced segregation, and the delayed conversion of bubbles to voids. For 20%-cold-worked type 316 stainless steel irradiated in the HFIR, helium is trapped on the high density of dislocations, leading to the nucleation of a high density of bubbles. Because of the large number of small bubbles present, they become the dominant sinks for both vacancies and interstitials, and swelling is inhibited. In solution-annealed steel, the opposite occurs. Helium is not effectively trapped, and bubble nucleation occurs on a coarser scale than in cold-worked material. This accelerates the transition of cavities from bubbles to voids and leads to greater swelling than that observed in the absence of high helium concentrations. The precipitation of fine titanium-rich carbides (designated as MC, where the M indicates the carbide contains more than one type of metal atom) in the PCA steel for fusion (Table 1 ) enhances bubble nucleation per increment of generated helium (Ref 4). As long as fine dispersions of MCs are preserved during irradiation, the association of fine MCs with helium bubbles hinders bubble coarsening by coalescence. The resulting high density of bubble/precipitate sinks also suppresses radiation-induced segregation and thus further enhances MC stability. This is illustrated in Fig. 3 , which compares the swelling behavior and micro-structure of cold-worked type 316 stainless steel and PCA steel. The low-swelling PCA steel contains a high density of small bubbles, compared to the large voids in the stainless steel. Helium-enhanced MC stability and suppressed radiation-induced segregation are essential in extending the transient regime of swelling for fusion compared to fast-breeder reactor irradiation. The performance of PCA steel illustrated in Fig. 3 is for an He:dpa ratio of approximately 60, while the value for a fusion reactor is expected to be about 10 to 12. It still must be determined whether the metal carbide will remain stable under fusion reactor conditions. Fig. 3 Comparison of the swelling behavior and microstructure of cold-worked type 316 stainless steel and cold-worked PCA steel irradiated in the HFIR at 600 °C (1110 °F). (a) Cavity volume swelling versus neutron fluence. (b) Microstructure of 316 stainless steel after about 43 dpa. (c) Microstructure of the PCA steel after about 43 dpa. Source: Ref 4


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Swelling in Ferritic Steels. When chromium-molybdenum ferritic steels such as 21=4Cr-1Mo, 9Cr-1MoVNb, and 12Cr-1MoVW (Table 1 ) were irradiated in fast reactors to neutron fluences of 17.6 × 1026 n/m2 and a displacement-damage level of up to 80 dpa (helium concentrations 10 appm He) (Ref 9). Furthermore, significantly greater amounts of precipitation occurred at 400 and 500 °C (750 and 930 °F) under conditions of increased void formation. It appeared that these precipitate changes were the result of irradiation-induced solute segregation, in which the migration of vacancies and helium to cavity surfaces is accompanied by the preferential migration of solute atoms away from these sites. This can result in irradiation-induced precipitate phases that are unstable under similar thermal conditions in the absence of neutron irradiation. Similar changes were noted for the 12Cr-1MoVW steel (Ref 9).

Mechanical Properties Although ferritic steels are more resistant to swelling than austenitic steels, irradiation may have a more critical effect on the mechanical properties of ferritic steels. In particular, the effect of irradiation on fracture behavior is of crucial importance in light-water reactors and may limit the use of ferritic steels in fusion reactors. As noted above, low-temperature irradiation can result in the formation of dislocation loops, solute clusters, vacancy clusters, precipitates, and microvoids. This microstructural alteration causes most of the changes in mechanical properties. Transmutation helium can also affect mechanical properties. Low-Temperature Tensile Behavior. An example of the effect of fast reactor irradiation on the strength of a ferritic steel is shown in Fig. 4 for 9Cr-1MoVNb steel irradiated in the EBR-II at 390 to 550 °C (735 to 1020 °F) and tested at the irradiation


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temperature. An increase in the yield stress is observed for irradiation up to about 425 °C (800 °F). No hardening occurred at 450 °C (840 °F) and above. With increasing fluence below approximately 425 °C (800 °F), strength increases to a saturation level, after which it remains unchanged. Depending on the steel and the fluence, irradiation-enhanced softening is possible at temperatures above about 425 °C (800 °F) because of recovery and precipitate coarsening, which can be hastened by the irradiation. The change in the ultimate tensile strength is similar to that of the yield stress. Changes in ductility reflect the strength changes; an increase in strength results in a ductility decrease and vice versa. In general, adequate ductility is maintained for these irradiation conditions. Similar observations have been made for 12Cr-1MoVW (Ref 11) and 21=4Cr-1Mo steels (Ref 12). Fig. 4 0.2% yield stress and ultimate tensile strength of 9Cr-1MoVNb steel as a function of test temperature for irradiated specimens (12 dpa), as-heat-treated controls, and thermally aged controls. The test temperature equals the irradiation and aging temperatures; specimens were aged 5000 h, which corresponded to the time in-reactor. Source Ref 10

To determine the effect of helium, nickel-doped 9Cr-1MoVNb and 12Cr-1MoVW steel tensile specimens were irradiated in the HFIR (Ref 13). The results were compared with results for undoped steels irradiated similarly and undoped and doped specimens irradiated in the EBR-II, in which little helium was generated. At 300 and 400 °C (570 and 750 °F), results indicated that the transmutation helium caused an increase in strength in addition to that caused by the displacement damage. No helium effect was apparent on specimens irradiated at 500 °C (930 °F). A qualitatively similar behavior is observed when austenitic stainless steels are irradiated in the EBR-II (little helium). For type 316 stainless steel, both solution-annealed (Ref 14) and 20%-cold-worked (Ref 15) steels have been investigated in some detail. An attempt has been made to model the behavior observed in the EBR-II in terms of the defects produced during irradiation (Ref 16). Only minor differences were observed between stainless steels irradiated in the EBR-II and the HFIR up to approximately 600 °C (1110 °F), indicating that helium had little effect on tensile behavior under these conditions (Ref 17). Elevated-Temperature Tensile Behavior⎯Helium Embrittlement. For elevated temperatures, displacement damage is no longer stable, and flow properties are basically unaffected by irradiation (Fig. 4 ). However, in certain irradiated alloys containing helium, the strength decreases upon irradiation at elevated temperatures, but the ductility also decreases (Ref 18). Total elongation measured in a tensile test drops to only a few tenths of a percent. Although a temperature of approximately 0.5 Tm is often associated with helium embrittlement, the temperature will depend on the helium concentration and the tensile strain rate. As the helium concentration increases and/or the strain rate decreases, the temperature at which helium embrittlement occurs will decrease. Elevated-temperature helium embrittlement is accompanied by intergranular fracture and is thought to be caused by helium on grain boundaries. For austenitic stainless steels, the effect can occur with the presence of only a few atomic parts per million of helium⎯even the small amounts formed during fast-reactor irradiation (Ref 18). Embrittlement is more severe for cold-worked than for solution-annealed material, although the effect generally appears at temperatures at which recrystallization or recovery of the cold-worked material occurs. The large difference between the embrittlement of cold-worked and solution-annealed austenitic steels may indicate that some grain-boundary migration (by recrystallization or grain growth) during irradiation or testing may be necessary to obtain the


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extremely low elongations found in austenitic stainless steels. The difference between the cold-worked and solution-annealed type 316 stainless steels might be explained by grain boundaries collecting helium during recrystallization of the cold-worked structure. Such embrittlement could impose an upper temperature limit on the use of austenitic steels. Again, alloying with titanium can improve resistance to embrittlement. The improvement is due to bubble refinement when helium is trapped on MC particles, as shown in Fig. 5 for PCA steel. The PCA-Al was in a solution-annealed condition with no MC particles (Fig. 5 c). The PCA-B3 was aged at 800 °C (1470 °F) prior to 25% cold working. When the helium was trapped in the fine bubbles formed on the metal carbides (Fig. 5 b), a relatively small change in ductility occurred upon irradiation at 600 °C (1110 °F) to approximately 22 dpa and 1750 appm of helium (Ref 17). Without helium trapping (Fig.c 5 ), the helium collected at grain boundaries, and a much larger decrease in ductility occurred. It should be noted, however, that irradiation to higher fluences will be required to determine if this resistance to helium embrittlement continues. Fig. 5 Tensile ductility (a) of solution-annealed PCA steel and aged and cold-worked PCA steel. Irradiation caused a large decrease in the ductility of the solution-annealed PCA steel but not the cold-worked steel. This difference was correlated with fine bubbles on the MC precipitates that were present in the aged and cold-worked steel (b) but not in the solution-annealed steel (c). Source: Ref 17


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All indications are that the martensitic steels, such as 9Cr-1MoVNb and 12Cr-1MoVW, are much more immune to helium embrittlement than the austenitic steels (Ref 19). The reasons for such resistance are not completely understood. Immunity is not inherent in the body-centered cubic (bcc) crystal structure (compared to the face-centered cubic structure), because helium embrittlement occurs in vanadium and niobium alloys. It appears likely that the resistance to helium embrittlement is related to the martensitic microstructure. In the normalized condition, martensite has a fine lath structure containing a high density of dislocations. After tempering, a ferrite matrix containing a high density of carbide particles and a lower dislocation density remains. However, the distinctive lath structure is still evident; long laths, separated mostly by low-angle boundaries, are grouped in packets. The packets and some laths are separated by high-angle boundaries. Prior-austenite grain boundaries are also present. This type of fine microstructure should allow the partitioning of helium atoms to the various boundaries, including the precipitate boundaries. Such a wide distribution of helium should effectively keep the helium concentration on a given high-angle grain boundary relatively low and should reduce the probability of intergranular failure (an intergranular failure is expected to propagate along high-angle boundaries).


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If the above conclusions are correct, then a much larger effect of helium would be expected on pure iron or on a steel with a polygonal ferrite structure (as opposed to a martensitic or bainitic microstructure). The many collection sites available in martensite are not present in pure iron, and the grain boundaries in such precipitate-free microstructures are more mobile at elevated temperatures. Indeed, this may explain the reduced ductility values observed on steels with a ferrite microstructure. For example, the total elongation of types 405 and 430 ferritic stainless steels containing 40 appm of helium decreased from 52 to 33% and 89 to 48%, respectively, when tested at 700 °C (1290 °F) (Ref 20). Irradiation Embrittlement. A Major concern for bcc ferritic steels involves the effect of irradiation on fracture toughness (Ref 21, 22). Irradiation can cause large increases in the ductile-to-brittle transition temperature (DBTT) and decreases in the upper-shelf energy (USE), as measured by Charpy V-notch specimens. Even if the DBTT of the unirradiated steel is below room temperature, it can be well above room temperature after irradiation. Irradiation embrittlement is related to the radiation-produced dislocation loops that form below 0.35 Tm; irradiation-induced precipitates can also have an effect. Loops are barriers to dislocation motion and give rise to the strengthening discussed above. The relationship of this increase in flow stress to irradiation embrittlement is shown schematically in Fig. 6 . Figure 6 shows how irradiation has shifted the flow stress upward. Under the assumption that the intersection of the fracture stress curve and the flow stress curve is the DBTT for the unflawed condition, the increase in flow stress is seen to cause a shift in the DBTT. Fig. 6 Schematic of suggested mechanism by which a strength increase due to irradiation causes an upward shift in the DBTT. Source: Ref 21

Although swelling and the other aspects of radiation damage do not play a role in light-water power reactors, irradiation embrittlement has been a major concern. Low-alloy pressure vessel steels specified by ASTM A 302-B and A 533-B are commonly used for this application. Shifts in DBTT of over 200 °C (360 °F) have been observed in A 302-B irradiated at less than 232 °C (450 °F) to fluences of approximately 1 × 1024 n/m2, which is less than 0.1 dpa (Fig. 7 ). Because displacement damage can be eliminated by annealing, the magnitude of the DBTT shift for a given fluence generally decreases with temperature. For a given temperature, embrittlement is rapid with increasing fluence at low fluence. A marked decrease in rate of embrittlement (as measured by the upward shift in DBTT) is then observed, and the rate appears to go to zero with increasing fluence (that is, saturation) (Fig. 7 ). Fig. 7 Effect of neutron fluence on the 41 J (30 ft · lbf) transition temperature in Charpy impact tests at temperatures below 232 °C (450 °F). Test specimen: 150 mm (6 in.) thick manganese-molybdenum steel (ASTM A 302, grade B). Source: Ref 21


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Irradiation embrittlement of the pressure vessel steels exposed to fluences typical of light-water reactors is affected by heat-to-heat variations, microstructure, and residual element content (Ref 21). Of the residual elements, copper and phosphorus have the greatest effect. One proposed mechanism for the effect of copper is that it enhances the formation of dislocation loops that lead to hardening (Ref 21). The effect of phosphorus has been attributed to a mechanism similar to the role it plays in temper embrittlement. This effect manifests itself under these conditions because of radiation-enhanced diffusion (Ref 21). Postirradiation heat treatment can restore the DBTT of low-alloy pressure vessel steels (Ref 21). Neutrons are moderated in light-water reactors to produce thermal neutrons, but a considerable flux of fast neutrons is still present. It is these neutrons with energies above about 1 MeV that are generally considered to produce the damage that causes the irradiation embrittlement (Ref 21). However, accelerated embrittlement was recently observed in surveillance specimens for the pressure vessel of the HFIR. These ASTM A 212-B steel specimens were irradiated at about 50 °C (120 °F) in a high thermal-to-fast-flux ratio position, where the measured property change was about an order of magnitude larger than that expected on the basis of the fast neutron fluence (Ref 23). Although thermal neutrons do not possess sufficient energy to dislodge an atom from the matrix, thermal neutrons can cause damage indirectly through transmutation reactions. In particular, a reaction between a thermal neutron and 56Fe to form 57Fe causes an atom recoil, when a γ-ray is released. Displacement damage by this recoil can cause the embrittlement observed in the HFIR (Ref 24). Recoil from a reaction between a thermal neutron and boron to form an α-particle can also cause displacement damage. Damage from these transmutation reactions becomes important whenever the displacement-damage energy deposited by these reactions comes within an order of magnitude of that deposited by fast neutrons (Ref 24). The displacement damage produced by thermal neutrons is believed to be more efficient than that produced by fast neutrons in causing microstructural changes leading to embrittlement. This is because the displacement cascades from the recoil reactions are smaller than those for fast neutrons; consequently, less in-cascade recombination (loss) of vacancies and interstitials takes place for thermal neutrons. Recoil displacement are expected to be especially important at low temperatures, as is the case in the support structure of light-water reactors (Ref 24). Irradiation embrittlement must also be considered in the development of ferritic steels for fast reactors and fusion reactors (Ref 22). The ferritic steels considered for use for fast-breeder reactors and fusion reactors are quite different from the low-alloy pressure vessel steels used for light-water reactors (Table 1 ). Furthermore, the types of fluences to which these steels will be exposed are considerably greater. While a light-water reactor steel will be irradiated to levels of the order of 1 × 1024 n/m2, which produces a displacement-damage level of less than 0.1 dpa, fluences of two to three orders of magnitude higher and damage levels exceeding 200 dpa are expected for fast-breeder and fusion reactors. Charpy impact tests for the pressure vessel steels are generally conducted using standard Charpy V-notch specimens. Because of space limitations in most test reactors used for the fast-breeder and fusion reactor programs, miniature Charpy V-notch (halfand third-size) specimens are used (Ref 25). Results from comparisons of the different specimen sizes have shown that the smaller specimens exhibit behavior relatively similar to that of large specimens (that is, increase in DBTT and decrease in USE). However, the magnitude of the USE is greatly reduced, and the DBTT is lower (Ref 25). Because of the appearance of a saturation in the shift in DBTT of the pressure vessel steels after an irradiation to approximately 2 × 1023 n/m2 (Fig. 7 ), miniature specimens of 12Cr-1MoVW steel were irradiated at 300 °C (570 °F) to typical light-water reactor fluences in the University of Buffalo reactor (UBR) to 8.6 × 1023 n/m2 (1 MeV) fluence below about 5 × 1020 n/cm2 (see the article "Corrosion in the Nuclear Power Industry" in Corrosion, Volume 13 of ASM Handbook, formerly 9th Edition Metals Handbook). High tensile stresses are not necessary for irradiation-assisted SCC, and cracks may occur at lower stresses for higher fluences. However, high tensile stress can exacerbate the problem. REFERENCES 1. R.L. Klueh and E.E. Bloom, Radiation Facilities for Fusion Reactor First Wall and Blanket Structural Materials Development, Nucl. Eng. Des., Vol 73, 1982, p 101−125 2. J.O. Stiegler and L.K. Mansur, Radiation Effects in Structural Materials, Ann. Rev. Mater, Sci., Vol 9, 1979, p 405−454 3. L.K. Mansur and E.E. Bloom, Radiation Effects in Reactor Structural Alloys, J. Met., Vol 34, 1982, p 23−31 4. P.J. Maziasz, Swelling and Swelling Resistance Possibilities of Austenitic Stainless Steel in Fusion Reactors, J. Nucl. Mater., Vol 122 and 123, 1984 p 472−486 5. F.A. Garner, Recent Insights on the Swelling and Creep of Irradiated Austenitic Alloys, J. Nucl. Mater., Vol 122 and 123, 1984, p 459−471 6. R.A. Weiner and A. Boltax, Comparison of High Fluence Swelling Behavior of Austenitic Stainless Steels, in Effects of Irradiation on Materials: Tenth Conference, STP 725, American Society for Testing and Materials, 1981, p 484−499 7. D.S. Gelles, Microstructural Examination of Several Commercial Ferritic Alloys Irradiated to High Fluence, J. Nucl. Mater., Vol 103 and 104, 1981, p 975−980 8. E.A. Little and D.A. Stowe, Void-Swelling in Irons and Ferritic Steels: II. An Experimental Survey of Materials Irradiated in a Fast Reactor, J. Nucl. Mater., Vol 87, 1979, p 25−39 9. P.J. Maziasz, R.L. Klueh, and J.M. Vitek, Helium Effects on Void Formation in 9Cr-1MoVNb and 12Cr-1MoVW Irradiated in HFIR, J. Nucl. Mater., Vol 141−143, 1986, p 929−937 10. R.L. Klueh and J.M. Vitek, Elevated-Temperature Tensile Properties of Irradiated 9Cr-1MoVNb Steel, J. Nucl. Mater., Vol 132, 1985, p 27−31 11. R.L. Klueh and J.M. Vitek, Tensile Behavior of Irradiated 12Cr-1MoVW Steel, J. Nucl. Mater., Vol 137, 1985, p 44−50 12. R.L. Klueh and J.M. Vitek, Tensile Properties of 21=4Cr-1Mo Steel Irradiated to 23 dpa at 390 to 550 °C, J. Nucl. Mater., Vol 140, 1986, p 140−148 13. R.L. Klueh and J.M. Vitek, Postirradiation Tensile Behavior of Nickel-Doped Ferritic Steels, J. Nucl. Mater., Vol 150, 1987, p 272−280 14. R.L. Fish and J.J. Holmes, Tensile Properties of Annealed Type 316 Stainless Steel After EBR-II Irradiation, J. Nucl, Mater.,


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Vol 46, 1973, p 113−120 15. R.L. Fish, N.S. Cannon, and G.L. Wire, Tensile Property Correlations for Highly Irradiated 20 Percent Cold-Worked Type 316 Stainless Steel, in Effects of Radiation on Structural Materials, STP 683, American Society for Testing and Materials, 1979, p 450−465 16. G.D. Johnson, F.A. Garner, H.R. Brager, and R.L. Fish. Microstructural Interpretation of the Fluence and Temperature Dependence of the Mechanical Properties of Irradiated AISI 316, in Effects of Radiation on Materials: Tenth Conference, STP 725, American Society for Testing and Materials, 1981, p 393−412 17. P.J. Maziasz, A Perspective on Present and Future Alloy Development Efforts on Austenitic Stainless Steels for Fusion Application, J. Nucl, Mater., Vol 133−134, 1985, p 134−140 18. E.E. Bloom, Irradiation Strengthening and Embrittlement, in Radiation Damage in Metals, American Society for Metals, 1976, p 295−329 19. R.L. Klueh and J.M. Vitek, The Resistance of 9Cr-1MoVNb and 12Cr-1MoVW Steels to Helium Embrittlement, J. Nucl, Mater., Vol 117, 1983, p 295−302 20. D. Kramer, A.G. Pard, and C.G. Rhodes, A Survey of Helium Embrittlement of Various Alloy Types, in Irradiation Embrittlement and Creep in Fuel Cladding and Core Components, British Nuclear Energy Society, 1972, p 109−115 21. J.R. Hawthorne, Irradiation Embrittlement, in Treatise on Materials Science and Technology, Vol 25, Academic Press, 1983, p 461−524 22. G.E. Lucas and D.S. Gelles, The Influence of Irradiation on Fracture and Impact Properties of Fusion Reactor Materials, J. Nucl. Mater., Vol 155−157, 1988, p 164−177 23. R.K. Nanstad, K. Farrell, D.N. Braski, and W.R. Corwin, Accelerated Neutron Embrittlement of Ferritic Steels at Low Fluence: Flux and Spectrum Effects, J. Nucl. Mater., Vol 158, 1988, p 1−6 24. L.K. Mansur and K. Farrell, On Mechanisms by Which a Soft Neutron Spectrum May Induce Accelerated Embrittlement, J. Nucl. Mater., to the published 25. W.R. Corwin and A.M. Hougland, Effect of Specimen Size and Material Condition on the Charpy Impact Properties of 9Cr-1Mo-V-Nb Steel, in The Use of Small-Scale Specimens for Testing Irradiated Material, STP 888, American Society for Testing and Materials, 1986, p 325 26. J.M. Vitek, W.R. Corwin, R.L. Klueh, and J.R. Hawthorne, On the Saturation of the DBTT Shift of Irradiated 12Cr-1MoVW With Increasing Fluence, J. Nucl. Mater., Vol 141−143, 1986, p 948−953 27. W.L. Hu and D.S. Gelles, The Ductile-to-Brittle Transition Behavior of Martensitic Steels Neutron Irradiated to 26 dpa, in Influence of Radiation of Material Properties: 13th International Symposium (Part II), STP 956, American Society for Testing and Materials, 1987, p 83−97 28. R.L. Klueh, J.M. Vitek, W.R. Corwin, and D.J. Alexander, Impact Behavior of 9-Cr and 12-Cr Ferritic Steels After Low-Temperature Irradiation, J. Nucl. Mater., Vol 155−157, 1988, p 973−977 29. W.R. Corwin, J.M. Vitek, and R.L. Klueh, Effect of Nickel Content of 9Cr-1MoVNb and 12Cr-1MoVW Steels on the Aging and Irradiation Response of Impact Properties, J. Nucl. Mater., Vol 149, 1987, p 312−320 30. D.S. Gelles and L.K. Thomas, Effects of Neutron Irradiation on Microstructure in Experimental and Commercial Ferritic Steels, in Ferritic Alloys for Use in Nuclear Energy Technologies, The Metallurgical Society, 1984, p 559−568 31. C. Wassilew, Influence of Helium Embrittlement on Post-Irradiation Creep Rupture Behavior of Austenitic and Martensitic Stainless Steels, in Nuclear Technology and Applications of Stainless Steels at Elevated Temperatures, The Metals Society, 1982, p 172−181 32. B. Van der Schaaf, The Effect of Neutron Irradiation on the Fatigue and Fatigue-Creep Behaviour of Structural Materials, J. Nucl. Mater., Vol 155−157, 1988, p 156−163 33. W.A. Coghlan, Recent Irradiation Creep Result, Int. Met. Rev., Vol 31, 1986, p 241−290 34. A.J. Jacobs and G.P. Wozadlo, Irradiation-Assisted Stress Corrosion Cracking as a Factor in Nuclear Power Plant Aging, in Proceeding of the International Conference on Nuclear Power Plant Aging, Availability Factor, and Reliability Analysis, American Society for Metals, 1985, p 173−180


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Low-Temperature Properties of Structural...

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Service Characteristics of Carbon and Low-Alloy Steels Low-Temperature Properties of Structural Steels Mamdouh M. Salama, Conoco Inc. CRITICAL STRUCTURAL COMPONENTS must be fabricated from steels that exhibit adequate low-temperature fracture toughness because of the serious consequences of failure due to brittle fracture. Codes used for the design of offshore structures specify low-temperature toughness requirements, and steel specifications that satisfy these requirements have been developed. The need for steels with higher fracture toughness and better weldability, as well as lower cost, has prompted major advancements in structural steel technology. These advancements are highlighted by the development of controlled-rolled and accelerated-cooled steels. This article reviews fracture resistance assessment procedures for welded joints and includes discussions on fatigue crack growth and fracture toughness. Fracture toughness requirements specified by different design codes are presented, and American Petroleum Institute (API), British Standards Institution (BSI), and American Society for Testing and Materials (ASTM) specifications for offshore structural steels are summarized, and applications of these specifications are discussed. This article also focuses on advances made in steel technology and the impact of these advances on the fracture toughness of steel.

Design and Failure Criteria Three major factors contribute to service failure of steel structures: • Brittle failure due to the presence of fabrication defects • Fatigue crack development • Crack development as a result of accidental damage It is not practical or economical to fabricate defect-free structures. Although the use of appropriate inspection and quality control procedures can limit the size of defects, it cannot eliminate defects entirely. Proper fatigue design practices and in-service inspection can control the growth of fatigue cracks, however, complete elimination of small fatigue cracks is unrealistic, particularly for complex welded structures. Ductile failure due to growth of fatigue cracks to a large, plastic collapse critical size is a rare event, but it is still more common than brittle fracture, especially in structures subjected to the turbulent North Sea environment. Ductile failure in the absence of cracklike defects is experienced only in cases of accidental overloads that grossly exceed normal design stresses. In addition to catastrophic failures of ships, tankers, offshore structures, pipelines, bridges, and vessels (Ref 1,2,3), numerous minor brittle failures of structures under construction or in service have resulted in delays and expensive repairs. To minimize the probability of these failures, the design of modern structures is based on the combined use of the methods of both classical design and structural integrity design. Structural integrity design is employed to prevent structural failure due to brittle fracture or premature fatigue cracking. Integrity design provides a tool for assessing fracture resistance by integrating stress analysis with evaluations of fabrication quality and the mechanical properties of the steel. The mechanical properties that are evaluated include fatigue crack growth, fracture toughness, and basic tensile properties (for example, yield strength and tensile strength). Currently, all design guidelines, codes, or standards for critical applications emphasize fracture control procedures that provide for the evaluation of properties such as fracture toughness, weldability, and strength. Stringent steel qualification criteria have contributed to the development of low-cost structural steels possessing superior mechanical properties. These structural steels combine desired properties such as higher strength, improved weldability, and higher fracture toughness in one product. These properties are vital in steels used for offshore structures because the inaccessibility of these structures makes in-service inspection and repair very difficult and extremely expensive.

Assessment of Fracture Resistance The offshore industry has used several advanced fracture mechanics methodologies and tests to establish allowable final defect, af (Ref 4). These include crack tip opening displacement (CTOD), tests, and, to a lesser extent, crack growth resistance,


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01 Sep 2005

JR, and failure assessment diagram methods. The CTOD approach allows calculation of size of the allowable final defect, af, using the following expression for the ratio between the critical defect and the CTOD value (Ref 5):

(Eq 1) where af is the half-length of a through-thickness rectilinear crack (for surface and buried cracks, and for crack geometries other than through-thickness rectilinear cracks, the equivalent through-thickness cracks, can be estimated by utilizing relationships available in fracture mechanics handbooks or from the literature (for example, Ref 6), Y/E is the ratio between the yield strength and the modulus of elasticity of the material, S/Y is the ratio between the nominal applied stress and the yield strength, SCF is the stress concentration factor, and α is the stress relief parameter, which equals 1.0 for no stress relief (that is, residual stress equals Y) and equals 0.0 for full stress relief (that is, no residual stress). Ensuring against brittle fracture by specifying a blanket CTOD value is difficult without performing detailed fatigue life calculations. However, toughness specifications in terms of CTOD values are valuable when used in conjunction with fatigue crack growth rate data in the framework of fracture mechanics analysis; they can provide useful information on tolerable defects, remaining product life, and allowable loading conditions. Because of the complexity of CTOD testing (Ref 7), most design codes still rely on Charpy V-notch (CVN) energy and transition temperature concepts as the main fracture toughness acceptance criteria. The CVN impact test is performed following international standards such as ASTM A 370 or BSI 131.

Fracture Toughness Requirements Almost all design guidelines for critical structures specify a minimum fracture toughness requirement. This section summarizes toughness requirements given by three existing design guidelines used in the offshore industry. These guidelines have been developed by API, the United Kingdom Department of Energy (DEn), and Det Norske Veritas (DNV), a Norwegian ship classifier. The fracture toughness criteria used in these guidelines are based mainly on CVN impact energy and transition temperature criteria. According to API RP 2A (Ref 8), underwater joints should meet notch toughness requirements as established by either the Naval Research Laboratory drop-weight test (ASTM E 208) or the CVN impact energy test. For the drop-weight test, the joints should be rated for no-break performance. The CVN test is performed on transverse test specimens. The minimum Charpy energy is specified as a function of the minimum yield strength of the steel (Table 1 ). The test temperature is specified as a function of the lowest anticipated service temperature (LAST) and the pipe diameter-to-thickness ratio, D/t: Test temperature(b)

Condition of testpiece

°C

°F

30 LAST − 20 LAST − 36 (a) D/t, diameter-to-thickness ratio. (b) LAST, lowest anticipated service temperature

Flat plate

D/t(a)

Table 1 Yield strength and impact energy guidelines for low-temperature structural steels Minimum Charpy V-notch impact energy Minimum yield strength

Average Transverse

Individual

Longitudinal

Transverse

Longitudinal

Specification

MPa

ksi

J

ft · lbf

J

ft · lbf

J

ft · lbf

J

ft · lbf

API RP 2A

52

48

35

...

...

41(a)

29.8(a)

...

...

Y(b)

Y(b)

0.10Y

0.10Y

...

...

0.07Y

0.07Y

...

...

Y ≤ 275

Y ≤ 40

18

13

27

20

14(c)

9.8(c)

20(c)

15(c)

United Kingdom DEn DNV

Y > 275 Y > 40 0.07 Y 0.07 Y 0.10 Y 0.10 Y 0.053 Y(c) 0.053 Y(c) 0.075 Y(c) 0.075 Y(c) (a) Minimum individual impact energy = minimum average −7 J, or −5.2 ft ·lbf. (b) Y, minimum yield strength of the thinnest plate. (c) Minimum individual CVN value = 0.75 × (minimum required average value)

For a D/t less than 20, the Charpy specimen is machined from as-fabricated pipe, and the test temperature is 10 °C (18 °F) below LAST. On the other hand, for a higher D/t, testing is conducted on samples machined from flat plates. A lower temperature


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01 Sep 2005

is specified for this test to account for possible toughness deterioration due to strains induced by forming (D/t of 20 is equivalent to 5% strain). No recommendations about the effect of plate thickness on toughness requirements are provided by the API specification. The DEn specification (Ref 9) provides fracture toughness criteria for structures located in the North Sea. The toughness criteria depend on the minimum yield strength and thickness of the plate, and on the postweld heat treatment, stress concentration, and location of test specimens. The recommended minimum average CVN values of transverse specimens are listed in Table 1 , and the recommended test temperatures are summarized in Table 2 . Table 2 Recommended CVN test temperatures according to United Kingdom Department of Energy specifications Test Temperature As-welded Thickness, t mm

in. 25

t≤20

t≤ =32

20 < t ≤ 100

25

Postweld heat treatment

Highly stressed

=32 < t ≤ 4

40 < t ≤ 100 19=16 < t ≤ 4

Others

Highly stressed

Others

Test location

°C

°F

°C

°F

°C

°F

°C

°F

Subsurface

−20

−5

−20

−5

−20

−5

−10

15

Subsurface

−40

−40

−30

−20

−30

−20

−20

−5

Mid-thickness

−30

−20

−20

−5

−20

−5

−20

−5

The DNV standard (Ref 10) provides fracture tougness requirements using both CVN and CTOD approaches. The recommended average minimum energy level depends on the yield strength as given in Table 1 . Table 3 lists the recommended impact-testing temperature in terms of the design temperature, TD, where TD is defined as 5 °C (9 °F) below the most probable lowest monthly mean temperature. The DNV specifications include minimum CTOD value requirements for weld procedure qualifications for plates 50 mm (2 in.) and greater in thickness. The requirements at the minimum design temperature are 0.35 mm (0.014 in.) for as-welded or local postweld heat-treated conditions, and 0.25 mm (0.010 in.) after furnace postweld heat treatment. Table 3 Recommended CVN test temperature according to DNV specifications Test temperature(a) Thickness, t

Special steel

mm

in.

Primary steel

Secondary steel

°C

°F

°C

°F

°C

°F

TD

TD

TD

TD

...

...

t ≤ 12.5

t ≤ =2

12.5 < t ≤ 25.5

1

=2 < t ≤ 1

TD − 20

TD − 36

TD

TD

...

...

25.5 < t ≤ 50

1 50

t>2

TD − 40

TD − 72

TD − 40

TD − 72

TD − 20

TD − 36

1

(a) TD, design temperature

There are some differences in toughness requirements among the design codes. Also, requirements continue to change with new issues of each code as more data become available. Table 4 highlights the importance of low-temperature toughness requirements for offshore structural steels by comparing the toughness requirements for two cases using the different codes. For thicker plates and for structures used in more severe environments such as the Arctic, higher fracture toughness values are required. Table 4 Low-temperature fracture toughness requirements for 50 mm (2 in.) thick plate used in offshore structures Minimum water temperature Structure location Gulf of Mexico

North Sea

Minimum yield strength

Minimum individual CVN impact energy

Temperature

°C

°F

Specification

MPa

ksi

J

ft · lbf

5

40

API

345

50

27

20

°C

°F

−25

−15

415

60

41

30

−25

−15

50

25

18

−40

−40

5

40

DEn

345 415

60

30

22

−40

−40

5

40

DNV

345

50

25

18

−35

−30

415

60

30

22

−40

−40

345

50

27

20

−30

−20

415

60

41

30

−30

−20

345

50

25

18

−40

−40

0 0

32 32


API DEn

Page 1016



0

32

DNV

01 Sep 2005

415

60

30

22

−40

−40

345

50

25

18

−40

−40

415

60

30

22

−45

−50

Fatigue Crack Growth in Structural Steel Steel structures generally include complex welded joints that have large local stress concentrations and that are subject to fatigue loadings induced by environmental forces. This, in addition to fabrication defects that are often present in welded structures, will result in the early initiation of fatigue cracks. Eventually the crack grows to a size at which failure may occur. Therefore, the fatigue life of welded components can be estimated by integrating an appropriate crack growth equation such as Paris law between the allowable initial defect, ai, and the final defect at which failure occurs, af. The size of the final defect depends on the fracture toughness of the material and the applied stress. The Paris equation (Eq 2 ) is bounded by the threshold value, ∆Ko, and the critical value, Kmax, which is a measure of the fracture toughness as shown in Fig. 1 . The crack growth equation provides a relationship between the crack growth rate, da/dN, and the stress intensity factor range, ∆K in the following form:

(Eq 2) The stress intensity factor range, ∆K, is defined by:

(Eq 3) where ∆S is the cyclic stress range, F is a correction factor dependent on component and crack geometries, and a is the half-length of a through-thickness rectilinear crack. Fig. 1 Idealized fatigue crack growth model

The crack growth parameters C and m are experimentally determined constants that depend on the material, loading condition, and environment. Reference 6 provides a C value of 3 × 10−13 and an m value of 3, in units of N (load) and mm (length), for ferritic steels with yield strengths up to 600 MPa (87 ksi). These values are based on the upper limit of air fatigue data shown in Fig. 2 for a variety of weld metals and heat-affected zone (HAZ) microstructures (Ref 11). Since the development of these data, extensive fatigue crack growth data have been developed for offshore structural steels such as BS 4360 grade 50D. Figure 3 presents crack growth data for this steel in both an air environment and in free corrosion and cathodic protection (CP) conditions in seawater. The CP levels were between −800 and −1100 mV (with respect to a silver/silver chloride reference electrode). Figure 3 was developed using data from research done exclusively on rectangular through-thickness notched parent steel specimens (no weld metal or HAZ data are included) (Ref 12). The test frequency varied between 0.1 and 1.0 Hz, the temperature between 5 and 20 °C (40 and 70 °F) and the stress ratio, R, between 0.0 and 0.5. Based on these data, a C value of 2.3 × 10−12 and an m value of 3 (in units of N and mm) have been suggested (Ref 13). These values predict crack growth rates that are about one order of magnitude higher than crack growth rates calculated using the values recommended in BS PD6493. Fatigue lives are therefore reduced by one order of magnitude due to the interactive effect of crack size. Fig. 2 Fatigue crack growth in the weld metal and heat-affected zones of carbon-manganese steel base plates in an air environment


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01 Sep 2005

Fig. 3 Fatigue crack growth of BS 4360 grade 50D steel in air and in free corrosion and cathodic protection conditions in seawater


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01 Sep 2005

p p The Paris crack growth equation is generally valid within the ∆K range of 300 to 1800 MPa mm) (9 to 52 ksi in:). Values p p mm (9 ksi of K below about 300 MPa p in:) fall in the threshold range where crack propagation does not occur, and values p above about 1800 MPa mm (52 ksi in:) fall in the range where the static mode of fracture occurs as the fracture toughness limit of the material is approached. The following relationship between ∆Ko and the applied stress ratio, R, is provided in BS PD6493:

(Eq 4) In Eq 4 , ∆Ko depends on R. The Paris equation (Eq 2 ) does not depend on R. The relationship presented in Eq 4 provides the lower bound to all published threshold data for grade 50D steel in air and seawater (Ref 9). It has been suggested that other data for similar steels and for austenitic steels lie below the PD6493 line (Ref 10). Including these data, the following relationship, based on a 97.7% probability of survival for the data in Fig. 4 , has been recommended (Ref 10):

(Eq 5) Fig. 4 Fatigue crack growth threshold data for ferritic steels with yield strengths up to 600 MPa (87 ksi)


Page 1019



01 Sep 2005

Steel Specifications Structural steel specifications are generally based on the appropriate national or industry standards such as ASTM, API, BSI, and so on. In most cases, standards provide mainly basic requirements such as limits on chemical composition and tensile properties. During the mid-1960s, several in-service and structural fabrication problems in Gulf of Mexico offshore structures were encountered. These problems indicated that common pipes such as API 5L B and structural steels such as ASTM A 7 and ASTM A 36 do not always meet the design or service needs of the offshore industry (Ref 3). Failure analysis studies on several salvaged structures have shown that low notch toughness, laminations in the steel, lamellar tearing, and poor weldability were major contributors to the failures. These results made offshore operators and certifying authorities cognizant of the need for more restrictive standards to ensure that the steels used for offshore applications are of high quality and satisfy strict fracture toughness and weldability requirements. Therefore, standards such as API 2H, 2Y, and 2W (Ref 14,15,16) were developed. The types of structural steels that are addressed in these standards include killed fine-grain normalized, controlled-rolled, and quenched and tempered steels, as well as controlled-rolled accelerated-cooled (referred to as thermomechanical control process, or TMCP), steel. In addition to the above API grades, special grades from general standards such as ASTM and BSI are also specified for offshore structures. Table 5 summarizes the chemical composition and mechanical properties of some offshore structural steels. There are several differences among these specifications in the details they provide regarding limitations on steelmaking, chemical composition, mechanical properties, and quality. Table 5 Chemical composition and mechanical properties of selected offshore structural steels Chemical composition, maximum wt%(a) Specific ation Grade API 2H

API 2W

Condition

C

Si

Mn

P

Nb

Al, total

Ni

Cr

42

Normalized

0. 0.15− 0.90− 0. 0.01 18 0.40 1.35 03 0

0.02− 0.01 2 0.06

...

...

...

...

...

289 (42)

427−56 5 (62−82)

50

Normalized

0. 0.15− 1.15− 0. 0.01 0.01−0.04 0.02− 0.01 18 0.40 1.60 03 0 2 0.06

...

...

...

...

...

345 (50)

483−62 0 (70−90)

42

TMCP

0. 0.15− 0.90− 0. 0.01 16 0.50 1.35 03 0

0.03

0.02− 0.01 0.75 0.25 0.08 0.35 0.003−0. 290−4 2 0.06 02 62 (42−6 7)

427 (62)

50

TMCP

0. 0.15− 1.15− 0. 0.01 16 0.50 1.60 03 0

0.03

0.02− 0.01 0.75 0.25 0.08 0.35 0.003−0. 345−4 2 0.06 02 83 (50−7 5)

448 (65)

60

TMCP

0. 0.15− 1.15− 0. 0.01 16 0.50 1.60 03 0

0.03

0.02− 0.01 1.0 0.25 0.15 0.35 0.003−0. 414−6 2 0.06 02 21

517 (75)


S

0.04

N

Mo Cu

Ti

Yield Tensile streng strengt th(b), h(b), MPa MPa (ksi) (ksi)

Page 1020



01 Sep 2005

(60−9 0) API 2Y

BS 4360(c)

ASTM A 633

ASTM A 131

42

Q&T

0. 0.15− 0.90− 0. 0.01 16 0.50 1.35 03 0

0.08

0.02− 0.01 0.75 0.25 0.08 0.35 0.003−0. 290−4 2 0.06 02 62 (42−6 7)

427 (62)

50

Q&T

0. 0.15− 1.15− 0. 0.01 16 0.50 1.60 03 0

0.08

0.02− 0.01 0.75 0.25 0.08 0.35 0.003−0. 345−5 2 0.06 02 17 (50−7 5)

448 (65)

60

Q&T

0. 0.15− 1.15− 0. 0.01 16 0.50 1.60 03 0

0.08

0.02− 0.01 1.0 0.25 0.15 0.35 0.003−0. 414−6 2 0.06 02 21 (60−9 0)

517 (75)

43D Normalized

0. 16

0.50

1.50

0. 0.04 0.003−0.1 04 0 0(d) 0

...

...

...

...

...

...

...

270 (39)

430−51 0 (62−74)

50D Normalized

0. 0.10− 18 0.50

1.50

0. 0.04 0.003−0.1 04 0 0(d) 0

...

...

...

...

...

...

...

345 (50)

490−62 0 (71−90)

55E

Normalized

0. 0.10− 22 0.60

1.60

0. 0.04 0.003−0.1 04 0 0(e) 0

...

...

...

...

...

...

...

430 (62)

550−70 0 (80−101 )

C

Normalized

0. 0.15− 1.15− 0. 0.05 0.01−0.05 20 0.50 1.50 04 0 0

...

...

...

...

...

...

...

345 (50)

485−62 0 (70−90)

D

Normalized

0. 0.15− 1.00− 0. 0.05 20 0.50 1.60 04 0 0

...

...

...

0.25 0.25 0.08 0.35

...

345 (50)

485−62 0 (70−90)

EH32 Normalized

0. 0.10− 0.90− 0. 0.04 18 0.50 1.60 04

0.05

...

...

0.40 0.25 0.08 0.35

...

315 (46)

470−58 5 (68−85)

EH36 Normalized

0. 0.10− 0.90− 0. 0.04 18 0.50 1.60 04

0.05

...

...

0.40 0.25 0.08 0.35

...

360 (51)

490−62 0 (71−90)

... 0.7 0.7− 0.60 0.15 1.00 ... 515 585 0. 0.40 0.40− 0. 0.02 0.02 min A Quenched (75) (85) 07 (class and 1.0 −0.9 −0.2 −1.3 0.70 02 5 5 3) precipitation 0 5 0 heat treated (a) Q&T, quenched and tempered. (b) Heat analysis. (c) Values are selected for a thickness of 25 mm (1 in.); values may be reduced as thickness increases. (d) Ref 17. (e) 0.003−0.10% V. (f) 0.003−0.20% V ASTM A 70

Tables 6 and 7 compare BSI and API specifications for tensile strength and toughness properties of similar grades of steel. In addition to the differences in toughness values, there are differences in how each specification handles the effect of thickness on the yield strength. Furthermore, API 2W and 2Y provide not only minimum yield and tensile strength limits, but also an upper limit on the yield strength. A limit on the maximum yield strength is very important to ensure a reasonable match between the strength of the weld metal and the base plate. In general, it is desirable to ensure that the weld metal strength is higher than the steel strength. Table 6 Tensile properties of offshore low-temperature structural steels Thickness, t Specification BS 4360

Grade 43D, 43E

mm

Minimum yield strength in.

MPa

ksi 41

Minimum tensile strength MPa

ksi

t ≤ 16

t ≤ 0.63

280

430−510

62−74

16 < t ≤ 40

0.63 < t ≤ 1.6

270

39

430−510

62−74

40 < t ≤ 63

1.6 < t ≤ 2.5

255

37

430−510

62−74

63 < t ≤ 100

2.5 < t ≤ 4.0

240

35

430−510

62−74


Page 1021


t ≤ 0.63

355

51

490−620

71−90

16 < t ≤ 40

0.63 < t ≤ 1.6

345

50

490−620

71−90

40 < t ≤ 63

1.6 < t ≤ 2.5

340

49

490−620

71−90

63 < t ≤ 100

2.5 < t ≤ 4.0

(a)

(a)

490−620

71−90

t ≤ 16

t ≤ 0.63

450

65

550−700

80−101

16 < t ≤ 40

0.63 < t ≤ 1.6

430

62

550−700

80−101

40 < t ≤ 63

1.6 < t ≤ 2.5

415

60

550−700

80−101

55E

63 < t ≤ 100

2.5 < t ≤ 4.0

400

58

550−700

80−101

55F

63 < t ≤ 100

2.5 < t ≤ 4.0

(a)

(a)

550−700

80−101

42

42

427−565

62−82

55E, 55F

50 API 2W, 2Y

01 Sep 2005

t ≤ 16

50D, 50E, 50F

API 2H


42 50 60

t ≤ 63

t ≤ 2.5

289

t > 63

t > 2.5

289

42

427−565

62−82

t ≤ 63

t ≤ 2.5

345

50

483−620

70−90

t > 63

t > 2.5

324

47

483−620

70−90

290−462

42−67

427

62 62

t ≤ 25

t ≤ 1.0

t > 25

t > 1.0

290−427

42−62

427

t ≤ 25

t ≤ 1.0

345−517

50−75

448

65

t > 25

t > 1.0

345−483

50−70

448

65

t ≤ 25

t ≤ 1.0

414−621

60−90

517

75

t > 25

t > 1.0

414−586

60−85

517

75

(a) By agreement

Table 7 Toughness requirements for offshore low-temperature structural steels Minimum average energy, CVN Specification BS 4360

Grade 43D, normalized 43E, normalized 50D, normalized 50E, normalized 50F, quenched and tempered 55E, normalized 55F, quenched and tempered

API 2H, 2W, 2Y

API 2W, 2Y

Test temperature

J

ft · lbf

°C

°F

41

30

−10

15

27

20

−20

−5

61

45

−20

−5

27

20

−50

−60

41

30

−20

−5

27

20

−30

−20

47

35

−30

−20

27

20

−50

−60

47

35

−30

−20

27

20

−50

−60

61

45

−20

−5

27

20

−50

−60

41

30

−40

−40

−60

−75

27

20

42

34(a)(b)

25(a)(b)

−40

−40

42, S-2(c)

34(a)(b)

25(a)(b)

−60

−75

50

41(a)(b)

30(a)(b)

−40

−40

50, S-2(c)

41(a)(b)

30(a)(b)

−60

−75

60

48(a)(b)

35(a)(b)

−40

−40

60, S-2(c) 48(a)(b) 35(a)(b) −60 −75 (a) API CVN tests use transverse specimens. (b) API provides supplement S-7 for CVN tests using specimens uniformly strained 5% and aged at 250 °C (480 °F) for 1 h. (c) S-2 supplementary requirement


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In addition to the CVN toughness requirements in Table 7 , API provides two supplements using different toughness criteria. The first supplement is for toughness based on crack tip opening displacement tests of weld HAZ. Tests are performed in accordance with Section 3 of API RP 2Z (Ref 18); the heat input is 1.5 to 5 kJ/mm (38 to 125 kJ/in.), and the preheat is 100 to 250 °C (210 to 480 °F). For thicknesses up to and including 75 mm (3 in.), the required CTOD value is 0.25 mm (0.010 in.) at −10 °C (15 °F). For thicknesses greater than 75 mm (3 in.), the required CTOD value is 0.38 mm (0.015 in.) at −10 °C (15 F). The second supplement is for toughness of plates using the drop-weight test. The test is done in accordance with ASTM E 208 using P-3 specimens. The acceptable criterion is no-break performance at −35 °C (−30 °F). Although standards for offshore structural steels are generally more restrictive than those used by other industries, they provide only minimum requirements for tensile properties, fracture toughness properties, control of chemical composition, and dimensional tolerances. Therefore, offshore operators often include additional requirements in the steel purchase specifications. These specifications generally include additional limitations on chemical composition along with requirements for higher toughness, a weldability evaluation, reduced tolerances, and an increased frequency of testing. Table 8 compares the chemical compositions of typical offshore structural steels and the composition allowed by the API 2H (Ref 14). In the typical steels, limits are imposed on more elements, and the maximum limits of carbon, sulfur phosphorus, and carbon equivalent are reduced. These restrictions are intended to ensure improved toughness and weldability. Table 8 API composition specification (product analysis) for offshore low-temperature structural steels compared with typical United States and foreign specifications Composition, %

Specification

C

Mn

P

S

Si

API 2H (1988), grade 50

0.22, 1.07− max 1.60

0.04 max

Typical United States mill

0.15

1.34

0.015 0.006

0.30

Typical foreign mill

0.12

1.44

0.009 0.001

0.38

Nb

Ni

Cr

Mo

V

Cu

As

Sn

Carbon equivalent, Sb maximum(a)

(b)

(b)

(b)

(b)

(b)

(b)

(b)

(b)

0.43

0.040 0.17 0.08 0.05 0.00 0.03 (c) 6 2 2

(c)

(c)

0.40

Al, total

0.015 0.13− 0.005 0.015 max 0.45 −0.05 −0.06 0.05

0.020 0.035 0.18 0.00 0.00 0.00 0.16 0.00 0.00 0.00 0.38 9 1 1 3 1 0 (a) International Institute of Welding carbon equivalent (CE): CE = %C + (%Cr + %Mo + %V)/5 + (%Ni + %Cu)/15. (b) Not specified. (c) Not reported

Advances in Steel Technology Many significant advances in steelmaking processes have been made by steel companies to meet the demand for high-quality lower-cost structural steels with higher strength, improved weldability, and increased fracture toughness. These advances include the close monitoring of the supply of desulfurized iron in the blast furnace, the widespread use of continuous casting of thick slab for rolling to plate, the introduction of vacuum arc degassing, vacuum degassing, and argon stirring and injection techniques, along with almost exclusive use of basic oxygen process steelmaking (Ref 19). These advances in the steelmaking process have resulted in major improvements in structural steels, including significant control of alloying elements (for example, carbon, manganese, niobium, vanadium, and aluminum), major reductions in impurities (for example, sulfur, phosphorus, and nitrogen), and improved uniformity of composition and properties. Recent advances in computer control and rolling capacity have allowed the development of a new class of high-strength low-alloy steels, namely TMCP steels. The TMCP involves both controlled rolling and controlled (accelerated) cooling to produce steels with a very fine grain size (ASTM 10 to 12). The main aim of TMCP is to increase strength and fracture toughness and improve weldability by reducing the carbon equivalent and controlling the chemical composition (additional information is available in the article "Weldability of Steels" in this Volume). The API 2W specification (Ref 16) covers TMCP steel plates with minimum yield strengths between 290 and 415 MPa (42 and 60 ksi). Strength in TMCP steels is maximized by reducing the ferrite grain size and increasing the volume fraction of the second phase. Accelerated cooling is used to achieve these effects. The influence of cooling rate on strength and toughness is shown in Fig. 5 . A variation in cooling rate can be expected between surface and mid-thickness regions of thick plates. The addition of small amounts of niobium is very effective in strengthening the steel without affecting toughness (Fig. 6 ). However, the addition of more than 0.04% Nb is not desirable because it can cause a reduction in toughness, particularly in the subcritically reheated grain-coarsened heat-affected zone. Fig. 5 Effect of cooling rate on selected properties of TMCP steels. (a) Strength. (b) Toughness. Source: Ref 20


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01 Sep 2005

Fig. 6 Effect of niobium content on selected properties of TMCP steels. (a) Strength. (b) Toughness. Source: Ref 21

Fracture Toughness Characteristics of Structural Steels Fracture toughness of steel has improved greatly as a result of advances in steel technology. Figure 7 compares CVN transition curves of old and new steels using transverse subsurface specimens. Steel A is a 60 mm (23=8in.) carbon-manganese steel that was used in 1975. It has a carbon level of 0.21%. Steel B is a modern 70 mm (23=4in.) thick normalized carbon-manganese steel with a carbon level of 0.114% and some microalloying (0.29% Ni, 0.025% Nb, and 0.022% Cu). Steel C is a modern 50 mm (2 in.) thick controlled-rolled and accelerated-cooled TMCP steel with a carbon level of 0.11% and some microalloying (0.23% Ni, 0.03% Nb, and 0.24% Cu). The yield strengths of steels A, B, and C are 355, 369, and 506 MPa (51, 54, and 73 ksi), respectively. Figure 7 shows the improved fracture toughness of modern steels as indicated by a decrease in the transition temperature and an increase in the upper-shelf energy. Fig. 7 Correlation of (a) Charpy V-notch impact energy and (b) crystallinity with nil-ductility transition temperature (NDTT) for three steels: A, 60 mm (23=8in.) thick old carbon-manganese steel (0.21% C) with a yield strength of 355 MPa (51 ksi); B, 70 mm (23=4in.) thick modern carbon-manganese steel (0.114C-0.29Ni-0.025Nb-0.022 Cu) with a yield strength


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of 369 MPa (54 ksi); and C, 50 mm (2 in.) thick TMCP steel (0.11C-0.23Ni-0.03Nb-0.24Cu) with a yield strength of 506 MPa (73 ksi). Source: Ref 22

Figure 7 also presents nil-ductility transition temperature (NDTT) results for the three steels as determined by the drop-weight test (ASTM E 208). The results show that the NDTT for each steel corresponds to a different location on the CVN curve. The old steel, A, has an NDTT of −30 °C (−20 °F), which corresponds to a CVN transition temperature at 45 J (33 ft · lbf) and 85% crystallinity. The NDTT of B, the modern normalized steel, is −40 °C (−40 °F), which corresponds to a CVN transition temperature at 235 J (173 ft · lbf) and 8% crystallinity. The NDTT of C, the TMCP steel, is −60 °C (−75 °F), which corresponds to a CVN transition temperature at 200 J (150 ft · lbf) and 15% crystallinity. Test data obtained for API 2W grade 50, an 89 mm (31=2in.) thick TMCP steel, indicate a similar relationship between the NDTT and the CVN transition curve (Ref 23). The API 2W grade 50 steel has an NDTT of −60 °C (−75 °F), which corresponds to CVN impact energy at NDTT of about 200 J (150 ft · lbf); the CVN energy at a CTOD value of 0.25 mm (0.010 in) is about 150 J (110 ft · lbf) (Fig. 8 ). Fig. 8 Correlation of (a) CVN impact energy and crack tip opening displacement and (b) crystallinity with nil-ducility transition temperature for API 2W grade 50, a TMCP steel with a thickness of 89 mm (31=2in.)


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These test results show that modern steels have higher fracture toughness and lower NDTT than older steels. However, the results also raise questions about the suitability of the current practice of assessing the transition temperature of modern steels based on a CVN energy level of 25 to 45 J (18 to 33 ft · lbf) as shown in Table 4 . For older steels, CVN energy of 45 J (33 ft · lbf) corresponds to a temperature that is about the same as the NDTT, while for modern steels it corresponds to a temperature that is about 40 to 60 °C (105 to 140 °F) lower than the NDTT. This difference in steel behavior needs to be addressed in both design codes and steel specifications.

Fracture Toughness of Welded Structures Assessing the fracture toughness of offshore structural steel involves evaluating not only base plate toughness but also HAZ and weld metal toughness. Although both HAZ toughness and weld metal toughness requirements are usually included in fabrication specifications, HAZ toughness requirements are sometimes specified in the steel purchase agreement. Because the small amount of material at the tip of the sharp fatigue crack can be examined with CTOD testing, a detailed evaluation of the toughness of the different HAZ regions is possible (Ref 24). This positional accuracy allows for the identification of isolated regions in the HAZ with toughness substantially lower than that of the bulk material. These local brittle zones (LBZs) have occurred during testing for crack extension at stress levels far lower than those at which cracks extended in the bulk material. The presence of LBZs is not a new problem, nor is it limited to modern steels. In most steels, LBZs are associated with the grain-coarsened regions of the heat-affected zone (GCHAZ). Figure 9 identifies the different HAZ regions in a single-bevel multipass weld. Figure 9 also presents a plan view of a polished section illustrating a method for calculating the length and the percent of the grain-coarsened (GC) regions. Evaluation of wide plate tests suggests that fractures are likely to initiate from the GC areas where the grain size is greater than 80 µm (0.0024 in.), or ASTM No. 4 (Ref 22, 25). Fig. 9 Regions of the heat-affected zone. (a) The HAZ regions in a single-bevel multipass weld. SCHAZ, subcritical heat-affected zone; ICHAZ, intercritical heat-affected zone; FGHAZ, fine-grain heat-affected zone; SRGCHAZ, subcritically reheated grain-coarsened heat-affected zone; IRGCHAZ, intercritically reheated grain-coarsened heat-affected zone. (b) Plan view of a polished weld section showing a method for calculating the length and the percent of the GCHAZ. GC, grain-coarsened


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01 Sep 2005

There are several reasons for the current interest in LBZs. The need to reduce costs results in optimized structures that have less redundancy and a large number of highly stressed joints. To reduce welding costs, narrow groove preparation is used that may result in an HAZ that is normal to the loading direction. Also, unlike normalized steel in which the HAZ yield strength is higher than that of the base plate, TMCP steel sometimes has an HAZ yield strength that is lower than that of both the weld metal and the base plate (Fig. 10 ). Softening of the HAZ can also be expected if thermal cutting is used during fabrication. This behavior can limit the application of thermal cutting to TMCP steels. Fig. 10 Heat-affected zone hardness of conventional (normalized) steel and TMCP steel

Figure 11 shows the effect of weld metal strength on the CVN toughness of the HAZ for a carbon-manganese steel using the same narrow groove welding procedures. The yield strength of the steel is 458 MPa (66 ksi). Similar results were obtained using CTOD tests; the CGHAZ toughness decreases (that is, transition temperature increases) with increasing weld metal strength. Therefore, a highly overmatched weld metal (that is, a high ratio of weld metal yield strength, Yw to base metal yield strength Yb) is not desirable. The increase in weld metal strength appears to cause an unfavorable deformation state (constraint) that reduces toughness and enhances brittle crack initiation and propagation of GC-HAZ. These results also suggest that care must be taken


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01 Sep 2005

when correlating HAZ toughness values based on simulated microstructure with values obtained from actual welded joints. The combination of lower structural redundancy, higher stresses, and the location of lower-strength heat-affected zones normal to the loading directions can result in situations where fatigue cracks can propagate through more GC regions, thus increasing the possibility of brittle fracture. Toughness values supplied by ASTM and API are established by testing under ideal laboratory conditions and may not reflect actual HAZ toughness values experienced in the field. Customers should expect toughness values to be lower than those specified in standards. Fig. 11 Plot of impact energy versus temperature to show the effect of weld metal matching (ratio of weld metal yield strength, Yw, to base plate yield strength, Yb) on CVN transition curves for theheat-affected zone. Source: Ref 26

Grain-coarsened regions will always exist in structural steels, but their size depends on both the steel and the welding process used. Therefore, it is necessary to identify an acceptable size for these regions. One study indicates that as long as the percentage of GC regions sampled by the crack front of CTOD specimens is less than 7%, no low toughness values can be measured (Fig. 12 ). Similar evaluations have shown that a major deterioration in CTOD values results only when the GC regions sampled by the crack fron exceed 10% (Ref 28). Fig. 12 Crack tip opening displacement versus the percent of grain-coarsened regions for several structural steels. Source: Ref 27


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01 Sep 2005

Steel companies are responding to industry concerns about LBZs by working on the development of LBZ-free steels through appropriate microalloying (Ref 29). The development of an LBZ-free low aluminum-boron steel has been reported (Ref 30). The low aluminum-boron medium nitrogen chemistry (0.009% A1, 0.0024% B, 0.0052% N, with no copper or niobium additions) was designed to nucleate the maximum amounts of boron nitride precipitates during the cooling of the initial thermal cycle in the GCHAZ. This is intended to promote the austenite-to-ferrite transformation and to prevent the bainite transformation. Figure 13 compares the LBZ-free steel with a conventional TMCP steel. The two steels are of the 415-MPa (60-ksi) yield strength class. The cumulative distribution of the critical heat-affected zone CTOD values of the two steels were compared using 27 mm (11=16in.) CTOD specimens. These specimens were machined from joints welded using submerged arc welding with a heat input of 5.0 kJ/mm (125 kJ/in.). Fig. 13 Heat-affected zone toughness of low aluminum-boron LBZ-free TMCP steel and conventional TMCP steel. Heat input using submerged arc welding in 5.0 kJ/mm (125 kJ/in.). Source: Ref 30

ACKNOWLEDGMENT The author wishes to express his thanks to the management of Conoco Inc. for permission to publish this article.


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REFERENCES 1. W.S. Pellini, Guidelines for Fracture-Safe and Fatigue-Reliable Design of Steel Structures, The Welding Institute, 1983 2. S.T. Rolfe and J.M. Barsom, Fracture and Fatigue Control in Structures⎯Applications of Fracture Mechanics, Prentice-Hall, 1977 3. M.L. Peterson, Steel Selection for Offshore Structures, J. Petrol. Technol., Vol 27, 1975, p 274−282 4. H.C. Rhee and M.M. Salama, Application of Fracture Mechanics Method to Offshore Structural Crack Instability Analysis, J. Ocean Eng. Technol., Vol 1, (No. 1), 1987, p 94−103 5. M.G. Dawes and M.S. Kamath, The CTOD Design Curve Approach to Crack Tolerance, in Proceedings of the Conference on Tolerance of Flaws in Pressurized Components, Institution of Mechanical and General Technician Engineers, 1978 6. "Guidance on Some Methods for the Derivation of Acceptance Levels for Defects in Fusion Welded Joints," PD6493, British Standards Institution, 1980 7. "Method for Crack Opening Displacement (COD) Testing," BS 5762, British Standards Institution, 1979 8. "API Recommended Practice for Planning, Designing and Constructing Fixed Offshore Platforms," RP 2A, American Petroleum Institute, 1989 9. "Offshore Installations: Guidance on Design and Construction," United Kingdom Department of Energy, 1985 10. "Rules for the Design, Construction and Inspection of Offshore Structures," Det Norske Varitas, 1977 11. S.J. Maddox, Fatigue Crack Propagation Data Obtained from Parent Plate, Weld Metal and HAZ in Structural Steels, Weld. Res. Int., Vol 4 (No. 1.), 1974 12. G.S. Booth and S.J. Dobbs, Corrosion Fatigue Crack Growth in BS 4360 Grade 50D Steel⎯An Analysis, Weld. Inst. Res. Bull., Vol 27 (No. 9), 1986, p 293−297 13. S.J. Maddox, Revision of the Fatigue Clauses in BS PD 6493, in Proceedings of the International Conference on Weld Failures, The Welding Institute, 1988, p P47-1 to P47-15 14. "Specification for Carbon Manganese Steel Plates for Offshore Platform Tubular Joints," API 2H, American Petroleum Institute, 1989 15. "Specifications for Steel Plates, Quenched-and-Tempered, for Offshore Structures," API 2Y, American Petroleum Institute, 1989 16. "Specification for Steel Plates for Offshore Structures, Produced by Thermo-mechanical Control Process (TMCP)," API 2W, American Petroleum Institute, 1989 17. "Specification for Weldable Structural Steels," BS 4360, British Standards Institution, 1979 18. "Recommended Practice for Preproduction Qualification for Steel Plates for Offshore Structures," RP 2Z, American Petroleum Institute, 1987 19. E.F. Walker, Steel Quality, Weldability and Toughness, in Steel in Marine Structures, C. Noordhoek and J. de Back, Ed., Elsevier Science Publishers, 1987, p 49−69 20. Y. Nakano, K. Amano, J, Kudo, E. Kobayashi, T. Ogawa, S. Kaihara, and A. Sato, "Preheat and PWHT-Free 150-mm Thick API 2W Grade 60 Steel Plate for Offshore Structures," in Proceedings of the 7th International Conference on Offshore Mechanics and Arctic Engineering, M.M. Salama et al., Ed., Vol 3, American Society of Mechanical Engineers, 1988, p 89−101 21. T. Shiwaku, T. Shimohata, S. Takashima, H, Kaji, and K. Masubuchi, YS 420 and 460 MPa Class High Strength Steel Plates for Arctic Offshore Structures Manufactured by TMCP, in Proceedings of the 7th International Conference on Offshore Mechanics and Arctic Engineering, M.M. Salama et al., Ed., Vol 3, American Society of Mechanical Engineers, 1988, p 95−101 22. A.C. de Koning, J.D. Harston, K.D. Nayler, and R.K. Ohm, Feeling Free Despite LBZ, in Proceedings of the 7th International Conference on Offshore Mechanics and Arctic Engineering, M.M. Salama et al., Ed., Vol 3, American Society of Mechanical Engineers, 1988, p 161−179 23. M. Kurihara, H. Kagawa, and I. Watanbe, Coarse Grain HAZ Toughness Evaluation on Heavy Gauge TMCP Steel Plate By Wide Plate Test, in Proceedings of the 8th International Conference on Offshore Mechanics and Arctic Engineering, M.M. Salama et al., Ed., Vol 3, American Society of Mechanical Engineers, 1989, p 649−656 24. S.E. Webster and E.F. Walker, The Significance of Local Brittle Zones to the Integrity of Large Welded Structures, in Proceedings of the 7th International Conference on Offshore Mechanics and Arctic Engineering, M.M. Salama et al., Ed., Vol 3, American Society of Mechanical Engineers, 1988, p 395−403 25. R.M. Denys, Fracture Control and Brittle Zones, A General Appraisal, in Proceedings of the International Conference on Weld Failures, The Welding Institute, 1988 p P44-1 to P44-17 26. M. Kocak, L. Chen, and G. Gnirss, Effects of Notch Position and Weld Metal Matching on CTOD of HAZ, in Proceedings of the International Conference on Weld Failures, The Welding Institute, 1988 p P7-1 to P7-10 27. D.P. Fairchild, Fracture-Toughness Testing of Weld Heat-Affected Zones in Structural Steel, in Fatigue and Fracture Testing of Weldments, STP 1058, H.I. McHenry and J.M. Potter, Ed., American Society for Testing and Materials, 1990


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28. K. Hirabayashi, H. Harasawa, H. Kobayashi, T. Sakaurai, M. Hano, and T. Yasuoka, Welding Procedures of Offshore Structure to Achieve Toughness of the Welded Joint, in Proceedings of the 6th International Conference on Offshore Mechanics and Arctic Engineering, M.M. Salama et al., Ed., Vol 3, American Society of Mechanical Engineers, 1987, p 151−157 29. K. Ohnishi, S. Suzuki, A. Inami, R. Someya, S. Sugisawa, and J. Furusawa, Advanced TMCP Steel Plates for Offshore Structures, in Microalloyed HSLA Steels: Proceedings of Microalloying 88, AMS INTERNATIONAL, 1988, p 215−224 30. S. Suzuki, K. Arimochi, J. Furusawa, K. Bessyo, and R. Someya, Development of LBZ-Free Low A1-B-Treated Steel Plates, in Proceedings of the 8th International Conference on Offshore Mechanics and Arctic Engineering, M.M. Salama et al., Ed., American Society of Mechanical Engineers, 1989, p 657−663


Page 1031


Fatigue Resistance of Steels

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Service Characteristics of Carbon and Low-Alloy Steels Fatigue Resistance of Steels Bruce Boardman, Deere and Company, Technical Center FATIGUE is the progressive, localized, and permanent structural change that occurs in a material subjected to repeated or fluctuating strains at nominal stresses that have maximum values less than (and often much less than) the tensile strength of the material. Fatigue may culminate into cracks and cause fracture after a sufficient number of fluctuations. The process of fatigue consists of three stages: • Initial fatigue damage leading to crack initiation • Crack propagation to some critical size (a size at which the remaining uncracked cross section of the part becomes too weak to carry the imposed loads) • Final, sudden fracture of the remaining cross section Fatigue damage is caused by the simultaneous action of cyclic stress, tensile stress, and plastic strain. If any one of these three is not present, a fatigue crack will not initiate and propagate. The plastic strain resulting from cyclic stress initiates the crack; the tensile stress promotes crack growth (propagation). Careful measurement of strain shows that microscopic plastic strains can be present at low levels of stress where the strain might otherwise appear to be totally elastic. Although compressive stresses will not cause fatigue, compressive loads may result in local tensile stresses. In the early literature, fatigue fractures were often attributed to crystallization because of their crystalline appearance. Because metals are crystalline solids, the use of the term crystallization in connection with fatigue is incorrect and should be avoided.

Fatigue Resistance Variations in mechanical properties, composition, microstructure, and macrostructure, along with their subsequent effects on fatigue life, have been studied extensively to aid in the appropriate selection of steel to meet specific end-use requirements. Studies have shown that the fatigue strength of steels is usually proportional to hardness and tensile strength; this generalization is not true, however, for high tensile strength values where toughness and critical flaw size may govern ultimate load carrying ability. Processing, fabrication, heat treatment, surface treatments, finishing, and service environments significantly influence the ultimate behavior of a metal subjected to cyclic stressing. Predicting the fatigue life of a metal part is complicated because materials are sensitive to small changes in loading conditions and stress concentrations and to other factors. The resistance of a metal structural member to fatigue is also affected by manufacturing procedures such as cold forming, welding, brazing, and plating and by surface conditions such as surface, roughness and residual stresses. Fatigue tests performed on small specimens are not sufficient for precisely establishing the fatigue life of a part. These tests are useful for rating the relative resistance of a material and the baseline properties of the material to cyclic stressing. The baseline properties must be combined with the load history of the part in a design analysis before a component life prediction can be made. In addition to material properties and loads, the design analysis must take into consideration the type of applied loading (uniaxial, bending, or torsional), loading pattern (either periodic loading at a constant or variable amplitude or random loading), magnitude of peak stresses, overall size of the part, fabrication method, surface roughness, presence of fretting or corroded surface, operating temperature and environment, and occurrence of service-induced imperfections. Traditionally, fatigue life has been expressed as the total number of stress cycles required for a fatigue crack to initiate and grow large enough to produce catastrophic failure, that is, separation into two pieces. In this article, fatigue data are expressed in terms of total life. For the small samples that are used in the laboratory to determine fatigue properties, this is generally the case; but, for real components, crack initiation may be as little as a few percent or the majority of the total component life. Fatigue data can also be expressed in terms of crack growth rate. In the past, it was commonly assumed that total fatigue life consisted mainly of crack initiation (stage I of fatigue crack development) and that the time required for a minute fatigue crack to grow and produce failure was a minor portion of the total life. However, as better methods of crack detection became available, it was discovered that cracks often develop early in the fatigue life of the material (after as little as 10% of total lifetime) and grow continuously until catastrophic failure occurs. This discovery has led to the use of crack growth rate, critical crack size, and


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fracture mechanics for the prediction of total life in some applications. Hertzberg's text (Ref 1) is a useful primer for the use of fracture mechanics methods.

Prevention of Fatigue Failure A thorough understanding of the factors that can cause a component to fail is essential before designing a part. Reference 2 provides numerous examples of these factors that cause fracture (including fatigue) and includes high-quality optical and electron micrographs to help explain factors. The incidence of fatigue failure can be considerably reduced by careful attention to design details and manufacturing processes. As long as the metal is sound and free from major flaws, a change in material composition is not as effective for achieving satisfactory fatigue life as is care taken in design, fabrication, and maintenance during service. The most effective and economical method of improving fatigue performance is improvement in design to: • • • • •

Eliminate or reduce stress raisers by streamlining the part Avoid sharp surface tears resulting from punching, stamping, shearing, and so on Prevent the development of surface discontinuities or decarburizing during processing or heat treatment Reduce or eliminate tensile residual stresses caused by manufacturing, heat treating, and welding Improve the details of fabrication and fastening procedures

Control of or protection against corrosion, erosion, chemical attack, or service-induced nicks and other gouges is an important part of proper maintenance of fatigue life during active service life. Reference 3 contains numerous papers pertaining to these subjects.

Symbols and Definitions In most laboratory fatigue testing, the specimen is loaded so that stress is cycled either between a maximum and a minimum tensile stress or between a maximum tensile stress and a specified level of compressive stress. The latter of the two, considered a negative tensile stress, is given an algebraic minus sign and called the minimum stress. Applied Stresses. The mean stress, Sm, is the algebraic average of the maximum stress and the minimum stress in one cycle:

(Eq 1) The range of stress, Sr, is the algebraic difference between the maximum stress and the minimum stress in one cycle: Sr = Smax − Smin (Eq 2) The stress amplitude, Sa, in one-half the range of stress:

(Eq 3) During a fatigue test, the stress cycle is usually maintained constant so that the applied stress conditions can be written Sm ± Sa, where Sm is the static or mean stress and Sa is the alternating stress equal to one-half the stress range. The positive sign is used to denote a tensile stress, and the negative sign denotes a compressive stress. Some of the possible combination of Sm and Sa are shown in Fig. 1 . When Sm = 0 (Fig. 1 a), the maximum tensile stress is equal to the maximum compressive stress; this is called an alternating stress, or a completely reversed stress. When Sm = Sa (Fig. 1 b), the minimum stress of the cycle is zero; this is called a pulsating, or repeated, tensile stress. Any other combination is known as an alternating stress, which may be an alternating tensile stress (Fig. 1 c), an alternating compressive stress, or a stress that alternates between a tensile and a compressive value (Fig. 1 d). Fig. 1 Types of fatigue test stress. (a) Alternating stress in which Sm = 0 and R = −1. (b) Pulsating tensile stress in which Sm = Sa, the minimum stress is zero, and R = 0. (c) Fluctuating tensile stress in which both the minimum and maximum stresses are tensile stresses and R = 1=3. (d) Fluctuating tensile-to-compressive stress in which the minimum stress is a compressive stress, the maximum stress is a tensile stress, and R = −1=3


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Nominal axial stresses can be calculated on the net section of a part (S = force per unit area) without consideration of variations in stress conditions caused by holes, grooves, fillets, and so on. Nominal stresses are frequently used in these calculations, although a closer estimate of actual stresses through the use of a stress concentration factor might be preferred. Stress ratio is the algebraic ratio of two specified stress values in a stress cycle. Two commonly used stress ratios are A, the ratio of the alternating stress amplitude to the mean stress (A = Sa/Sm) and R, the ratio of the minimum stress to the maximum stress (R= Smin/Smax). The five conditions that R can take range from +1 to −1: • • • • •

Stresses are fully reversed: R = −1 Stresses are partially revered: R is between −1 and zero Stress is cycled between a maximum stress and no load: The stress ratio R becomes zero Stress is cycled between two tensile stresses: The stress ratio R becomes a positive number less than 1 An R stress ratio of 1 indicates no variation in stress, and the test becomes a sustained-load creep test rather than a fatigue test

S-N Curves. The results of fatigue tests are usually plotted as the maximum stress or stress amplitude versus the number of cycles, N, to fracture, using a logarithmic scale for the number of cycles. Stress may be plotted on either a linear or a logarithmic scale. The resulting curve of data points is called an S-N curve. A family of S-N curves for a material tested at various stress ratios is shown in Fig. 2 . It should be noted that the fully reversed condition, R = −1, is the most severe, with the least fatigue life. For carbon and low-alloy steels, S-N curves (plotted as linear stress versus log life) typically have a fairly straight slanting portion


Page 1034



01 Sep 2005

with a negative slope at low cycles, which changes with a sharp transition into a straight, horizontal line at higher cycles. Fig. 2 Best-fit S-N curves for unnotched 300M alloy forging with an ultimate tensile strength of 1930 MPa (280 ksi). Stresses are based on net section. Testing was performed in the transverse direction with a theoretical stress concentration factor, Kt, of 1.0. Source: Ref 4

An S-N curve usually represents the median, or B50, life, which represents the number of cycles when half the specimens fail at a given stress level. The scatter of fatigue lives covers a very wide range and can occur for many reasons other than material variability. A constant-lifetime diagram (Fig. 3 ) is a summary graph prepared from a group of S-N curves of a material; each S-N curve is obtained at a different stress ratio. The diagram shows the relationship between the alternating stress amplitude and the mean stress and the relationship between maximum stress and minimum stress of the stress cycle for various constant lifetimes. Although this technique has received considerable use, it is now out of date. Earlier editions of the Military Standardization Handbook (Ref 5) used constant lifetime diagrams extensively, but more recent editions (Ref 4) no longer include them. Fig. 3 Constant-lifetime fatigue diagram for AISI-SAE 4340 alloy steel bars, hardened and tempered to a tensile strength of 1035 MPa (150 ksi) and tested at various temperatures. Solid lines represent data obtained from unnotched specimens; dashed lines represent data from specimens having notches with Kt = 3.3. All lines represent lifetimes of ten million cycles. Source: Ref 5

Fatigue limit (or endurance limit) is the value of the stress below which a material can presumably endure an infinite number


Page 1035



01 Sep 2005

of stress cycles, that is, the stress at which the S-N diagram becomes and appears to remain horizontal. The existence of a fatigue limit is typical for carbon and low-alloy steels. For many variable-amplitude loading conditions this is true; but for conditions involving periodic overstrains, as is typical for many actual components, large changes in the long-life fatigue resistance can occur (see the discussion in the section "Comparison of Fatigue Testing Techniques" in this article). Fatigue strength, which should not be confused with fatigue limit, is the stress to which the material can be subjected for a specified number of cycles. The term fatigue strength is used for materials, such as most nonferrous metals, that do not exhibit well-defined fatigue limits. It is also used to described the fatigue behavior of carbon and low-alloy steels at stresses greater than the fatigue limit. Stress Concentration Factor. Concentrated stress in a metal is evidenced by surface discontinuities such as notches, holes, and scratches and by changes in microstructure such as inclusions and thermal heat affected zones. The theoretical stress concentration factor, Kt, is the ratio of the greatest elastically calculated stress in the region of the notch (or other stress concentrator) to the corresponding nominal stress. For the determination of Kt, the greatest stress in the region of the notch is calculated from the theory of elasticity or by finite-element analysis. Equivalent values may be derived experimentally. An experimental stress concentration factor is a ratio of stress in a notched specimen to the stress in a smooth (unnotched) specimen. Fatigue notch factor, Kf, is the ratio of the fatigue strength of a smooth (unnotched) specimen to the fatigue strength of a notched specimen at the same number of cycles. The fatigue notch factor will vary with the life on the S-N curve and with the mean stress. At high stress levels and short cycles, the factor is usually less than at lower stress levels and longer cycles because of a reduction of the notch effect by plastic deformation. Fatigue notch sensitivity, q, is determined by comparing the fatigue notch factor, Kf, and the theoretical stress concentration factor, Kt, for a specimen of a given size containing a stress concentrator of a given shape and size. A common definition of fatigue notch sensitivity is:

(Eq 4) in which q may vary between 0 (where Kf = 1, no effect) and 1 (where Kf = Kt, full effect). This value may be stated as a percentage. As the fatigue notch factor varies with the position on the S-N curve, so does notch sensitivity. Most metals tend to become more notch sensitive at low stresses and long cycles. If they do not, it may be that the fatigue strengths for the smooth (unnotched) specimens are lower than they could be because of surface imperfections. Most metals are not fully notch sensitive under high stresses and a low number of cycles. Under these conditions, the actual peak stress at the base of the notch is partly in the plastic strain condition. This results in the actual peak stress being lower than the theoretical peak elastic stress used in the calculation of these theoretical stress concentration factor.

Stress-Based Approach To Fatigue The design of a machine element that will be subjected to cyclic loading can be approached by adjusting the configuration of the part so that the calculated stresses fall safely below the required line on an S-N plot. In a stress-based analysis, the material is assumed to deform in a nominally elastic manner, and local plastic strains are neglected. To the extent that these approximations are valid, the stress-based approach is useful. These assumptions imply that all the stresses will essentially be elastic. The S-N plot shown in Fig. 4 presents data for AISI-SAE 4340 steel, heat treated to a tensile strength of 1035 MPa (150 ksi) in the notched and unnotched condition. Figure 5 shows the combinations of cyclic stresses that can be tolerated by the same steel when the specimens are heat treated to different tensile strengths ranging from 860 to 1790 MPa (125 to 260 ksi). Fig. 4 Room temperature S-N curves for notched and unnotched AISI 4340 alloy steel with a tensile strength of 860 MPa (125 ksi). Stress ratio, R, equals −1.0. Source: Ref 4


Page 1036



01 Sep 2005

Fig. 5 Room temperature S-N curves for AISI 4340 alloy steel with various ultimate tensile strengths and with R = −1.0. Source: Ref 4

The effect of elevated temperature on the fatigue behavior of 4340 steel heat treated to 1035 MPa (150 ksi) is shown in Fig. 6 . An increase in temperature reduces the fatigue strength of the steel and is most deleterious for those applications in which the stress ratio, R, lies between 0.4 and 1.0 (Fig. 3 ). A decrease in temperature may increase the fatigue limit of steel; however, parts with preexisting cracks may also show decreased total life as temperature is lowered, because of accompanying reductions in critical crack size and fracture toughness. Fig. 6 S-N curves at various temperatures for AISI 4340 alloy steel with an ultimate tensile strength of 1090 MPa (158 ksi). Stress ratio, R, equals −1.0. Sources: Ref 4


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01 Sep 2005

Figure 7 shows the effect of notches on the fatigue behavior of the ultrahigh-strength 300M steel. A Kt, value of 2 is obtained in a specimen having a notch radius of about 1 mm (0.040 in.). For small parts, such a radius is often considered large enough to negate the stress concentration associated with any change in section. The significant effect of notches, even those with low stress concentration factors, on the fatigue behavior of this steel is apparent. Fig. 7 Room-Temperature S-N curves for a 300M steel with an ultimate tensile strength of 2000 MPa (290 ksi) having various notch severities. Stress ratio, R, equals 1.0. Source: Ref 4

Data such as those presented in Fig. 3 , 4 , 5 , 6 , and 7 may not be directly applicable to the design of structures because these graphs do not take into account the effect of the specific stress concentration associated with reentrant corners, notches, holes, joints, rough surfaces, and other similar conditions present in fabricated parts. The localized high stresses induced in fabricated parts by stress raisers are of much greater importance for cyclic loading than for static loading. Stress raisers reduce the fatigue life significantly below those predicted by the direct comparison of the smooth specimen fatigue strength with the nominal calculated stresses for the parts in question. Fabricated parts in simulated service have been found to fail at less than 50,000 repetitions of load, even though the nominal stress was far below that which could be repeated many millions of times on a smooth, machined specimen. Correction Factors for Test Data. The available fatigue data normally are for a specific type of loading, specimen size, and surface roughness. For instance, the R.R. Moore rotating-beam fatigue test machine uses a 7.5 mm (0.3 in.) diam specimen that is free of any stress concentrations (because of specimen shape and a surface that has been polished to a mirror finish), and that is subjected to completely reversed bending stresses. For the fatigue limits used in design calculations, Juvinall (Ref 6)


Page 1038



01 Sep 2005

suggests the correction of fatigue life data by multiplying the fatigue limit from testing, Ni, by three factors that take into account the variation in the type of loading, part diameter, and surface roughness: Design fatigue limit = Kl · Kd · Ks · Ni (Eq 5) where Kl is the correction factor for the type of loading, Kd for the part diameter, and Ks for the surface roughness. Values of these factors are given in Table 1 and Fig. 8 . Table 1 Correction factors for surface roughness (Ks), type of loading (Kl), and part diameter (Kd), for fatigue life of steel parts Value for loading in Factor

Bending

Torsion

Tension

1.0

0.58

0.9(a)

Ãwhere d ≤ 10 mm (0.4 in.)

1.0

1.0

1.0

Ãwhere 10 mm (0.4 in.) < d ≤ 50 mm (2 in.)

0.9

0.9

1.0

Kl Kd

Ks

See Fig. 8 .

(a) A lower value (0.06 to 0.85) may be used to take into account known or suspected undetermined bending because of load eccentricity. Source: Ref 6

Fig. 8 Surface roughness correction factors for standard rotating-beam fatigue life testing of steel parts. See Table 1 for correction factors from part diameter and type of loading. Source: Ref 6

Strain-Based Approach To Fatigue A strain-based approach to fatigue, developed for the analysis of low-cycle fatigue data, has proved to be useful for analyzing long-life fatigue data as well. The approach can take into account both elastic and plastic responses to applied loadings. The data are presented on a log-log plot similar in shape to an S-N curve; the value plotted on the abscissa is the number of strain reversals (twice the number of cycles) to failure, and the ordinate is the strain amplitude (half the strain range). During cyclic loading, the stress-strain relationship can usually be described by a loop, such as that shown in Fig. 9 . For purely elastic loading, the loop becomes a straight line whose slope is the elastic modulus, E, of the material. The occurrence of a hysteresis loop is most common. The definitions of the plastic strain range, ∆εp, the elastic strain range, ∆εc, the total strain range, ∆εt, and the stress range, ∆σ, are indicated in Fig. 9 . A series of fatigue tests, each having a different total strain range, will


Page 1039



01 Sep 2005

generate a series of hysteresis loops. For each set of conditions, a characteristic number of strain reversals is necessary to cause failure. Fig. 9 Stress-strain hysteresis loop. Source: Ref 7

As shown in Fig. 10 , a plot on logarithmic coordinates of the plastic portion of the strain amplitude (half the plastic strain range) versus the fatigue life often yields a straight line, described by the equation:

(Eq 6) where ε′f is the fatigue ductility coefficient, c is the fatigue ductility exponent, and Nf is the number of cycles to failure. Fig. 10 Ductility versus fatigue life for annealed AISI-SAE 4340 steel. Source: Ref 8

Because the conditions under which elastic strains have the greatest impact on fatigue behavior are the long-life conditions where stress-based analysis of fatigue is appropriate, the effects of elastic strain on fatigue are charted by plotting stress amplitude (half the stress range) versus fatigue life on logarithmic coordinates. As shown in Fig. 11 , the result is a straight line having the equation:

(Eq 7) where σ′f is the fatigue strength coefficient and b is the fatigue strength exponent.


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01 Sep 2005

Fig. 11 Strength versus fatigue life for annealed AISI-SAE 4340 steel. The equation for the actual stress amplitude, σa, is shown in ksi units. Source: Ref 8

The elastic strain range is obtained by dividing Eq 7 by E:

(Eq 8) The total strain range is the sum of the elastic and plastic components, obtained by adding Eq 6 and 8 (see Fig. 12 ):

(Eq 9) For low-cycle fatigue conditions (frequently fewer than about 1000 cycles to failure), the first term of Eq 9 is much larger than the second; thus, analysis and design under such conditions must use the strain-based approach. For long-life fatigue conditions (frequently more than about 10,000 cycles to failure), the second term dominates, and the fatigue behavior is adequately described by Eq 7 . Thus, it becomes possible to use Eq 7 in stress-based analysis and design. Fig. 12 Total strain versus fatigue life for annealed AISI-SAE 4340 steel. Data are same as in Fig. 10 and 11 . Source: Ref 8

Figure 13 shows the fatigue life behavior of two high-strength plate steels for which extensive fatigue data exist. ASTM A


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01 Sep 2005

440 has a yield strength of about 345 MPa (50 ksi); the other steel is a proprietary grade hardened and tempered to a yield strength of about 750 MPa (110 ksi). Under long-life fatigue conditions, the higher-strength steel can accommodate higher strain amplitudes for any specified number of cycles; such strains are elastic. Thus, stress and strain are proportional, and it is apparent that the higher-strength steel has a higher fatigue limit. With low-cycle fatigue conditions, however, the more ductile lower-strength steel can accommodate higher strain amplitudes. For low-cycle fatigue conditions (in which the yield strength of the material is exceeded on every cycle), the lower-strength steel can accommodate more strain reversals before failure for a specified strain amplitude. For strain amplitudes of 0.003 to 0.01, the two steels have the same fatigue life, 104 to 105 cycles. For this particular strain amplitude, most steels have the same fatigue life, regardless of their strength levels. Heat treating a steel to different hardness levels does not appreciably change the fatigue life for this strain amplitude (Fig. 14 ). Fig. 13 Total strain versus fatigue life for two high-strength low-alloy (HSLA) steels. Steels are ASTM A 440 having a yield strength of about 345 MPa (50 ksi) and a proprietary quenched and tempered HSLA steel having a yield strength of about 750 MPa (110 ksi). Source: Ref 7

Fig. 14 Effect of hardness level on plot of total strain versus fatigue life. These are predicted plots for typical medium-carbon steel at the indicated hardness levels. The prediction methodology is described under the heading "Notches." in this article.

Fuchs and Stephens's text (Ref 9), Proceedings of the SAE Fatigue Conference (Ref 10), and the recently published update to the SAE Fatigue Design Handbook (Ref 11) provide much additional detail on the use of state-of-the-art fatigue analysis methods. In fact, the chapter outline for the latter work, shown in Fig. 15 , provides an excellent checklist of factors to include in a fatigue analysis.


Page 1042



01 Sep 2005

Fig. 15 Checklist of factors in fatigue analysis. Source: Ref 11

Metallurgical Variables of Fatigue Behavior The metallurgical variables having the most pronounced effects on the fatigue behavior of carbon and low-alloy steels are strength level, ductility, cleanliness of the steel, residual stresses, surface conditions, and aggressive environments. At least partly because of the characteristic scatter of fatigue testing data, it is difficult to distinguish the direct effects of other variables such as composition on fatigue from their effects on the strength level of steel. Reference 3 addresses some excellent research in the area of microstructure and its effect on fatigue. Strength Level. For most steels with hardnesses below 400 HB (not including precipitation hardening steels), the fatigue limit is about half the ultimate tensile strength. Thus, any heat treatment or alloying addition that increases the strength (or hardness) of a steel can be expected to increase its fatigue limit as shown in Fig. 5 for a low-alloy steel (AISI 4340) and in Fig. 16 for various other low-alloy steels as a function of hardness. However, as shown in Fig. 14 for medium-carbon steel, a higher hardness (or strength) may not be associated with improved fatigue behavior in a low-cycle regime (45.7 m/min, or 150 sfm) are higher than those used in production practice for machining aircraft parts. At a constant metal removal rate that corresponds to a cutting speed of 53.3 m/min (175 sfm), ASP 60 and ASP 30 lasted 8 times and 4.5 times longer, respectively, than the M42 end mill. Other materials machined by P/M high-speed tool steel milling cutters include tough, hardened steels, such as 4340; austenitic stainless steels, such as AISI type 316; and nickel-base superalloys, such as Nimonic 80. Fig. 12 Results of end mill tests on Ti-6Al-4V. Hardness: 34 HRC


Page 1213


P/M Tool Steels

01 Sep 2005

Cutter, mm (in.)

25 (1) diam end mills

Feed, mm/tooth (in./tooth)

0.203 (0.008)

Radial depth of cut, mm (in.)

6.35 (0.250)

Axial depth of cut, mm (in.)

25.4 (1.000)

Cutting fluid

Soluble oil (1:20)

Tool life end point, mm (in.)

0.5 (0.020) wear

Hole Machining. Reamers, taps, and drills (Fig. 13 ) are also made from P/M high-speed tool steels. In one application, the tool life of 1=4-20 GH3 four-flute plug taps made from CPM Rex M4 and conventional M1, M7, and M42 were compared. The operation consisted of tapping a reamed 5.18 mm (0.204 in.) diam, 12.7 mm (0.500 in.) deep through hole in AISI 52100 steel at 32 to 34 HRC using a speed of 7.9 m/min (26 sfm) and chlorinated tapping oil. Eight to thirteen taps of each grade were tested. The CPM Rex M4 taps had an average tool life of 157 holes tapped before tool failure, compared to 35 holes for M1, 18 holes for M7, and 32 holes for M42. The tool life of the CPM Rex M4 in this application was about five times the life of conventional M42. Fig. 13 Reamers, taps, and drills made from P/M high-speed tool steels. Courtesy of Crucible Materials Corporation

Broaching tools constitute another major application for P/M high-speed tool steels because tool life is often improved when P/M steels are used to broach difficult-to-cut materials, such as case-hardened steels and superalloys. One application required broaching six ball tracks that are used in front-wheel-drive automobiles in constant velocity joint hubs made of a case-hardened steel. Figure 14 shows the joint hubs and the broaching tool used. In this broaching application,surface finish and form tolerance requirements are high because subsequent machining is not performed on the ball tracks. In broaching with low-carbon M35 tools with an 18.0 mm (0.709 in.) diameter and a 185.0 mm (7.283 in.) length, the total number of hubs machined per tool was 5600. The M35 tools experienced severe flank wear and developed a large built-up edge, which produced poor surface finishes. With an ASP 30 tool, 20,000 parts were produced.


Page 1214


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01 Sep 2005

Fig. 14 Broaching application. (a) Tool made from P/M high-speed tool steel that was used to produce ball tracks on joint hub. (b) ASP 30 tools produced 20,000 parts compared to 5600 parts by tools made from conventional high-speed tool steel. Courtesy of Speedsteel Inc.

Large broaching tools, such as those shown in Fig. 15 , are also being made from P/M high-speed tool steels, such as CPM M3 and M4, to upgrade the broach material and its performance. In general, large rounds for broaches are not available in conventional high-speed tool steels in sizes above about 254 mm (10 in.), but the larger sizes are available in P/M high-speed tool steels. One application for these tools is the broaching of involute splines in bores of truck transmission gear blanks. Bores up to 305 mm (12 in.) in diameter and 1380 mm (541=4in.) in length have been cut using such tools. Fig. 15 Large broaching tool made from P/M high-speed tool steel that was used for broaching involute splines in bores of truck transmission gear blanks. Courtesy of Crucible Materials Corporation

Gear Manufacturing. Gear hobs (Fig. 16 ) made from P/M high-speed tool steels can also provide substantial cost reductions by increasing machining rates. One application called for the hobbing of rear axle gears for heavy-duty trucks and tractor differentials. Hobs made of conventionally processed AISI M35 (65 HRC) and ASP 30 (67 HRC) were compared. Production results showed that the flank wear land on the hobs made of ASP 30 (0.44 mm, or 0.017 in.) was much less than on hobs made of M35 (0.71 mm, or 0.028 in.). Chipping of the edges was infrequent on the ASP 30 hobs, while the M35 hobs frequently displayed such damage. Fig. 16 Gear hobs made from P/M high-speed tool steels. Courtesy of Speedsteel Inc.


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01 Sep 2005

Sintered Tooling Figure 17 shows a number of P/M high-speed steel tools made by vacuum sintering cold isostatically pressed or mechanically pressed powders. Due to the higher temperatures used in processing, the primary carbides tend to be larger in vacuum-sintered, rather than hot isostatically pressed, high-speed steels. However, they are still smaller than in conventional high-speed steels of similar composition. Applications in which fully dense vacuum-sintered high-speed steels are currently in use include screw machine tooling, gear cutting tools, and indexable cutting tool inserts. Fig. 17 Examples of vacuum-sintered parts manufactured by the FULDENS process. Note the complexity of shapes attainable by this process. (a) Using cold isostatic pressing. (b) Using mechanical pressing

P/M Cold-Work Tool Steels A number of improved, high-vanadium P/M tool steels designed for high-wear and cold-work applications are commercially available. The chemical compositions of four representative P/M cold-work grades are given in Table 1 . As with P/M high-speed tool steels, the more uniform microstructure of the P/M cold-work steels yields better toughness. This is very important in


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P/M Tool Steels

01 Sep 2005

cold-work tooling because it allows higher hardnesses to be used with associated improvements in yield strength and wear resistance. Further, the use of higher vanadium contents in P/M cold-work tool steels than is practical in conventional cold-work tool steels has made for substantial improvements in wear resistance. Both CPM 9V and CPM 10V are examples of two P/M cold-work tool steels with outstanding wear resistance and toughness (Ref 1). CPM 9V is capable of being heat treated to 58 to 60 HRC using austenitizing temperatures at or above 1149 °C (2100 °F). However, much better toughness is obtained in the hardness range of 46 to 55 HRC using austenitizing temperatures ranging from 1038 to 1121 °C (1900 to 2050 °F). The impact toughness and wear resistance values of CPM 9V, compared to those of a number of conventional and P/M hot- and cold-work tool steels, are shown in Fig. 18 and 19 , respectively. In Fig. 19 , the wear resistance of CPM 9V at 53 to 55 HRC, which is the maximum hardness level recommended for good toughness, is clearly superior to that of conventional D2 (62 HRC) and CPM M4 (64 HRC). Fig. 18 Charpy C-notch impact properties of CPM 9V and other P/M and conventional tool steels at indicated hardnesses for cold-work applications

Fig. 19 Wear resistance of CPM 9V and other P/M and conventional tool steels at indicated hardnesses

CPM 10V is favored over CPM 9V in cold-work applications where greater hardness, higher compressive strength, or better wear resistance is required. In practice, CPM 10V has proved to be more wear resistant than any commercially available high-alloy tool steel, including the most highly alloyed high-speed tool steels. The outstanding wear resistance properties of CPM 10V are illustrated in Fig. 20 , which shows its laboratory wear test data compared to data for conventional D2, M2, and for CPM M4 at the maximum hardness levels typically used with these materials in cold-work tooling. Also included are wear data for D7, A7, and CPM T15. Although hardness has a major effect on wear resistance, its importance in this comparison is overshadowed by the very high vanadium carbide content of CPM 10V. Fig. 20 Wear resistance of CPM 10V and other P/M and conventional tool steels at indicated hardnesses


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01 Sep 2005

Figure 21 shows impact toughness data for CPM 10V compared to D2, M2, conventional M4, and CPM M4 at the same hardness levels used to generate the wear test data in Fig. 20 . These results, confirmed by actual field experience, indicate that CPM 10V exhibits toughness comparable to that of conventional D2 and M4. Fig. 21 Charpy C-notch impact properties of CPM 10V and other P/M and conventional tool steels at indicated hardnesses

Typical applications for CPM 10V include punches and dies for cold-forming and stamping operations, powder metal compaction tooling, roll-forming rolls, wood-working tools, wear parts, and so on. CPM 10V should also be considered as a cost-effective replacement for tungsten carbide and the Ferro-Tic steel-bonded carbides, particularly where these materials are prone to chipping and breakage or where the cost of these materials is prohibitive. Some examples of tooling and wear part applications where CPM 10V has demonstrated its cost effectiveness follow. Figure 22 shows a typical high-production-rate, progressive stamping die application in which CPM 10V punches were compared to D2 and CPM Rex M4. The material being stamped was a 0.381 mm (0.015 in.) thick copper-beryllium strip. In one operation, D2 piercing punches at 60 to 62 HRC averaged 75,000 parts before losing size; CPM Rex M4 at 64 HRC showed some signs of wear after 200,000 parts; and CPM 10V showed no wear after 400,000 parts and ultimately produced over 1,500,000 parts. In a coining operation in the same progressive die, CPM 10V was compared to D2 to produce a sharp radius bend indentation. D2 at 60 to 62 HRC required regrinding after 50,000 to 60,000 parts because of rounding of the punch nose radius, while the CPM 10V punch showed no signs of wear after 200,000 parts. Fig. 22 CPM 10V punch and copper-beryllium blank used in a progressive stamping operation. Courtesy of Crucible Materials Corporation


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Figure 23 shows a unique wear part application in which CPM 10V was selected both for its high wear resistance and its ability to be manufactured economically into the component shown. This part is an automotive push-rod assembly for diesel engines. CPM 10V was selected to eliminate a localized wear problem on the ball ends. Manufacturing of the balls requires cold or warm upsetting of 6.35 mm (0.250 in.) diam slugs cut from coil stock to form approximately 9.53 mm (0.375 in.) diam balls that are subsequently rough ground, drilled, heat treated, finish ground, and resistance welded to the push-rod stems. This part could not be produced economically in carbide or other pressed and sintered materials. Fig. 23 Automotive push-rod assembly with wear-resistant spherical CPM 10V tips. Courtesy of Crucible Materials Corporation

CPM 440V is a high-vanadium, high-chromium tool steel for applications requiring both high wear resistance and good corrosion resistance. The composition of this material (Table 1 ) is essentially that of T440C martensitic stainless steel to which about 5.75% V and increased carbon have been added to improve wear resistance. The wear resistance and toughness properties of CPM 440V are compared with those of T440C, D2, and CPM 10V in Table 7 . Although not as wear resistant as CPM 10V, CPM 440V exhibits outstanding wear resistance characteristics compared to conventional T440C and D2. Table 7 Comparative properties of CPM 440V, CPM 10V, conventional T440C, and D2 tool steel Crossed-cylinder wear resistance Alloy grade

Hardness

107 MPa

CPM 440V

59

276


Charpy C-notch toughness

1010 psi

J

ft · lbf

40

16.3

12

Page 1219


CPM 440V T440C

P/M Tool Steels

01 Sep 2005

56

...

...

21.7

16

56.5

28

4

35.3

26

D2

59

28

4

31.2

23

CPM 10V

60

517

75

35.3

26

P/M Hot-Work Tool Steels The absence of segregation in P/M tool steels makes them very attractive for hot-work tool and die applications because a frequent cause of premature failure of large die casting dies is thermal fatigue attributed to segregation and heterogeneous microstructures (Ref 20). As an example, H13 shot sleeves used to die cast four-cylinder engine blocks had an average service life of about 27,000 pieces when the starting material had little segregation, compared to about 13,000 pieces for material that showed clearly visible segregation in the starting material (Ref 21). During the past 10 years, efforts have been made to develop premium-quality, low-segregation tool and die steels by improved steelmaking processes such as vacuum degassing and vacuum arc or electroslag remelting. Powder metallurgy processing offers an alternative method of producing segregation-free hot-work tool steels of both standard and improved compositions and offers near-net shape capability. The compositions of three P/M hot-work tool steels now commercially available are given in Table 1 . Two are P/M versions of standard H13 and H19 hot-work tool steels, which, when made by P/M methods, have more uniform properties and equivalent or better toughness. The third is a high-vanadium modification of standard H19 with improved toughness and wear resistance. Figure 24 compares the longitudinal microstructures of P/M H13, standard conventional H13, and premium-quality conventional H13. Both of the conventional H13 alloys show significant microbanding, or alloy segregation, while the P/M H13 material has a homogeneous, segregation-free structure. Fig. 24 Longitudinal microstructure of (a) standard conventional H13, (b) premium-quality conventional H13, and (c) P/M H13. Vilella's etch. 50×

Results of property determinations made on conventional and P/M H13 to compare their heat treatment response, size change, toughness, tensile strength, and thermal fatigue resistance characteristics are given in Tables 8 , 9 , 10 , and 11 . The hardness characteristics across large sections of conventional and P/M H13 after a standard heat treatment [that is, 1010 °C (1850 °F)/1 h, air cool, 593 °C (1100 °F)/2 + 2 h] are given in Table 8 . These results show that the attainable hardnesses were the same but that the variation of hardness values across the P/M product was less than that of the conventional material, indicating greater uniformity. Table 9 gives the size change values for these products for the same heat-treated conditions. These results clearly show the isotropic nature of the P/M product. Table 8 Hardness of conventional and P/M H13 Section size in.

Hardness taken at 1.6 mm (1=16 in.) intervals across section, HRC

P/M H13

152

6 round

47.5−48.1

H13

127

5 round

46.0−47.7

Product(a)

mm

H13 152 × 406 6 × 16 46.0−48.6 (a) Preheated at 816 °C (1500 °F); austenitized at 1010 °C (1850 °F) for 1 h, air cooled; tempered at 593 °C (1100 °F) for 2 + 2 h


Page 1220


P/M Tool Steels

01 Sep 2005

Table 9 Size change of conventional and P/M H13 after heat treatment Section size Product(a)

mm

Size change, 0.0001 in./in. in.

Specimen location

Longitudinal direction

Transverse width

Transverse thickness

+6

+6

+6

P/M H13

152

6 round

Edge center

+6

+6

+6

H13

127

5 round

Mid-radius

−5

+8

+14

H13

152 × 406

6 × 16

Edge center

−1

+4

+7

+6 −2 (a) Preheated at 816 °C (1500 °F); austenitized at 1010 °C (1850 °F) for 1 h, air cooled; tempered at 593 °C (1100 °F) for 2 + 2 h

+17

Table 10 Impact and tensile properties of conventional and P/M H13

in.

Specimen location(b )

6 round

Section size Product(a)

mm

Charpy V-notch impact strength

0.2% yield strength

Tensile strength Elongation, Reduction MPa ksi % of area, %

J

ft · lbf

MPa

ksi

A

13.6

10

1407

204

1682

244

11

42

B

13.6

10

1407

204

1682

244

11

42

Tested at 21 °C (70 °F) P/M H13 H13

152 152 × 406

6 × 16

C

13.6

10

1413

205

1696

246

12

43

D

4

3

1400

203

1669

242

9

24

E

12.2

9

...

...

...

...

...

...

24.4

18

945

137

1213

176

17

51

Tested at 538 °C (1000 °F) P/M H13

152

6 round

A B

24.4

18

945

137

1213

176

17

51

H13

152 × 406

6 × 16

C

24.4

18

910

132

1200

174

17

53

D

16.3

12

924

134

1213

176

14

42

E 40.7 30 ... ... ... ... ... ... (a) Preheated at 816 °C (1500 °F); austenitized at 1010 °C (1850 °F) for 1 h, air cooled; tempered at 593 °C (1100 °F) for 2 + 2 h. (b) A, longitudinal; B, transverse; C, longitudinal center; D, transverse edge; E, longitudinal edge

Table 11 Thermal fatigue resistance of conventional and P/M H13 Section size Product(a) P/M H13

mm 152

in.

Average cycles to crack initiation

6 round

9000

H13 6000 152 × 406 6 × 16 (a) Preheated at 816 °C (1500 °F); austenitized at 1010 °C (1850 °F) for 1 h, air cooled; tempered at 593 °C (1100 °F) for 2 + 2 h

Table 10 gives the standard Charpy V-notch impact strength and tensile strength values obtained at 21 and 538 °C (70 and 1000 °F). At both test temperatures, the uniform properties of the P/M product are shown by the equal values obtained for longitudinal and transverse specimens and for edge and center locations. Compared to the conventional H13, the P/M product has slightly lower longitudinal toughness but higher transverse toughness. The tension test results show that the P/M and conventional products have equal tensile properties. Table 11 lists thermal fatigue test results obtained by alternately immersing specimens in molten lead at 621 °C (1150 °F) and water at 93 °C (200 °F) at a frequency of three cycles per minute. The average number of cycles to crack initiation was 50% greater for P/M H13. Initial field trials of CPM H19V have been encouraging. In one application in which the material was used to punch holes in forged hammer heads at elevated temperatures, it yielded 1700 to 4500 pieces, while conventional H13 averaged only 900 pieces before failure. Another major benefit of using the P/M process is the ability to produce near-net shape die cavities directly during HIP and thus minimize material input and subsequent machining. Figure 25 is an example of a die produced by ceramic core technology (Ref 22). In this process, a ceramic core material is cast in a reusable aluminum mold to the shape desired in the die cavity to be made. The ceramic core is then assembled with die steel powder in an outer steel container, as shown schematically in Fig. 26 . The dies can be made in pairs in a single outer container by using a steel separating plate. After powder loading, the assembly is evacuated, sealed, and hot isostatically pressed to compact the powder to full density and form the powder around the ceramic


Page 1221


P/M Tool Steels

01 Sep 2005

core. Typical HIP parameters for hot-work die steels are 1149 °C (2100 °F) and 103 MPa (15 ksi). Following HIP, the ceramic cores can be removed by grit blasting and/or chemical leaching. The die is then ready for heat treatment and finish machining operations. Fig. 25 P/M H19 die produced using prealloyed powder and the ceramic core process. Courtesy of Crucible Materials Corporation

Fig. 26 Schematic diagram of an assembly for producing P/M dies by the ceramic core process

The same process can be used to place cooling water passages within the die. An example is shown in Fig. 27 . With the P/M process, these passages can be designed and placed in a manner to make cooling most effective, rather than being restricted by machining limitations. With more effective cooling, die life would be expected to be enhanced. Fig. 27 P/M H13 die cut open to show cooling passages that were placed in the die using the ceramic core process. Courtesy of Crucible Materials Corporation


Page 1222


P/M Tool Steels

01 Sep 2005

Composite near-net shape dies can also be made by the P/M process. A highly alloyed material with high resistance to thermal fatigue or wear can be used in the die cavity while another lower cost or higher strength alloy is used in the remainder of the die. Figure 28 shows a P/M composite that uses Stellite 6 in the die cavity and H13 as the support material. Fig. 28 Composite P/M die made of Stellite 6 and H13 tool steel. (a) About 6.5% Stellite. (b) About 38% Stellite. Source: Ref 4

REFERENCES 1. R. Dixon, "Advances in the Development of Wear-Resistance High-Vanadium Tool Steels for Both Tooling and Non-Tooling Applications," Paper 8201-085, presented at the ASM Metals Congress, St. Louis, American Society for Metals, 1982 2. E. Bayer, HIP Tool Materials, Powder Metall., Vol 16 (No. 3), 1984, p 117−120 3. W. Stasko, V.K. Chandhok, and K.E. Pinnow, "Tool and Die Materials From Rapidly Solidified Powders," in Rapidly Solidified Materials: Properties and Processing, ASM INTERNATIONAL, 1988, p 49−57 4. H. Seilstorfer, PM Hot Work Tool Steels, Metal, Vol 42 (No. 2), Feb 1988, p 146−152 (in German) 5. V. Arnhold et. al., Cutting Tools From P/M High Speed Steels, Powder Metall. Int., Vol 21 (No. 2), 1989, p 67−74 (in German) 6. R. Riedl et al., Developments in High Speed Tool Steels, Steel Res., Vol 58 (No. 8), 1987, p 339−352 7. P. Beiss, PM Methods for the Production of High Speed Steels, Met. Powder Rep., Vol 38 (No. 4), April 1983, p 185−194 8. S. Karagoz and H. Fischmeister, Microstructure and Toughness in High Speed Tool Steels: The Influence of Hot Reduction and Austenitization Temperature, Steel Res., Vol 58 (No. 8), 1987, p 360 9. H. Berns et al., The Fatigue Behavior of Conventional and Powder Metallurgical High Speed Steels, Powder Metall. Int., Vol 19 (No. 4), 1987, p 22−26 10. N. Kawai and H. Takigawa, Methods for Producing PM High-Speed Steels, Met. Powder Rep., Vol 37 (No. 5), May 1982, p 237−240 11. P. Hellman et al., The ASEA-Stora Process, in Modern Developments in Powder Metallurgy, Vol 4, Plenum Press, 1970, p 573−582 12. E.J. Dulis and T.A. Neumeyer, Particle Metallurgy of High Speed Tool Steel, in Materials for Metal Cutting, Publication


Page 1223


P/M Tool Steels

01 Sep 2005

126, The Iron and Steel Institute, 1970, p 112−118 13. W.B. Kent, An Alternative Method of Processing High Speed Powder, in Processing and Properties of High Speed Tool Steels, Conference Proceedings, American Institute of Mining, Metallurgical, and Petroleum Engineers, 1980 14. M. Goransson et al., Method of Special Steel Production Via the STAMP Process, Met. Powder Rep., Vol 38 (No. 4), April 1983, p 205−209 15. B.A. Rickenson et al., CSD: New Horizons for Special Steels, in Towards Improved Performance of Tool Materials, The Metals Society, 1982, p 73−77 16. R.G. Brooks et al., The Osprey Process, Powder Metall., Vol 20, 1977, p 100−102 17. J. Smart et al., Pressing and Sintering Methods to Produce High Grade Tool Steels, Met. Powder Rep., Vol 35 (No. 6), June 1980, p 241−244 18. M.T. Podob and R.P. Harvey, Advantages and Applications of CMI's FULDENS Process, in Processing and Properties of High Speed Tool Steels, The Metallurgical Society, 1980, p 181−195 19. F.R. Dax, W.T. Haswell, and W. Stasko, Cobalt-Free CPM High Speed Steels, in Processing and Properties of High Speed Tool Steels, The Metallurgical Society, 1980, p 148−158 20. K.E. Thelning, "How Far Can H-13 Die Casting Die Steel Be Improved," Paper 114, presented at the Sixth SDCE International Die Casting Congress, Cleveland, OH, Society of Die Casting Engineers, 1970 21. J.H. Stuhl and A.M. Schindler, "New Materials Study of 5% Chromium Type Steels for Use in Die Casting Dies," Paper G-T75-053, presented at the Eighth SDCE International Die Casting Congress, Detroit, MI, Society of Die Casting Engineers, 1975 22. V.K. Chandhok, J.H. Moll, and M.E. Ulitchny, Process for Producing Parts With Deep Pocketed Cavities Using P/M Shape Technology, in Titanium Net Shape Technologies, F.H. Froes and D. Eylon, Ed., The Metallurgical Society, 1980 SELECTED REFERENCES • J.T. Berry, High Performance High Hardness High Speed Steels, Climax Molybdenum Company, 1970 • E.A. Carlson, J.E. Hansen, and J.C. Lynn, Characteristics of Full-Density P/M Tool Steel and Stainless Steel Parts, in Modern Developments in Powder Metallurgy, Vol 13, Metal Powder Industries Federation, 1980 • B.-A. Cehlin, "Improving Productivity With High Strength P/M High Speed Steel Cutting Tools," Paper MR82-948, presented at the Increasing Productivity With Advanced Machining Concepts Clinic, Los Angeles, CA, Society of Manufacturing Engineers, 1982 • P Hellman, Wear Mechanism and Cutting Performance of Conventional and High-Strength P/M High-Speed Steels, Powder Metall., Vol 25 (No. 2), 1982 • G. Hoyle, High Speed Steels, Butterworths, 1988 • A. Kasak and E.J. Dulis, Powder-Metallurgy Tool Steels, Powder Metall., Vol 21 (No. 2), 1978, p 114−123 • A. Kasak, G. Steven, and T.A. Neumeyer, High-Speed Tool Steels by Particle Metallurgy, SAE Paper 720182, Society of Automotive Engineers, 1972, p 2−5 • O. Siegwarth, Higher Productivity With ASP Tooling Material, Technical Paper MF81-137, Society of Manufacturing Engineers, p 1−22


Page 1224


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01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Specialty Steels and Heat-Resistant Alloys Maraging Steels Revised by Kurt Rohrbach and Michael Schmidt, Carpenter Technology Corporation MARAGING STEELS comprise a special class of high-strength steels that differ from conventional steels in that they are hardened by a metallurgical reaction that does not involve carbon. Instead, these steels are strengthened by the precipitation of intermetallic compounds at temperatures of about 480 °C (900 °F). The term maraging is derived from martensite age hardening and denotes the age hardening of a low-carbon, iron-nickel lath martensite matrix. Commercial maraging steels are designed to provide specific levels of yield strength from 1030 to 2420 MPa (150 to 350 ksi). Some experimental maraging steels have yield strengths as high as 3450 MPa (500 ksi). These steels typically have very high nickel, cobalt, and molybdenum contents and very low carbon contents. Carbon, in fact, is an impurity in these steels and is kept as low as commercially feasible in order to minimize the formation of titanium carbide (TiC), which can adversely affect strength, ductility, and toughness. Other varieties of maraging steel have been developed for special applications. Maraging steels are commercially produced by various steel companies in the United States and abroad. The absence of carbon and the use of intermetallic precipitation to achieve hardening produce several unique characteristics that set maraging steels apart from conventional steels. Hardenability is of no concern. The low-carbon martensite formed after annealing is relatively soft⎯about 30 to 35 HRC. During age hardening, there are only very slight dimensional changes. Therefore, fairly intricate shapes can be machined in the soft condition and then hardened with a minimum of distortion. Weldability is excellent. Fracture toughness is considerably better than that of conventional high-strength steels. This characteristic in particular has led to the use of maraging steels in many demanding applications.

Physical Metallurgy The unique characteristics of maraging steels have generated considerable interest, and extensive efforts have been devoted to gaining an understanding of the metallurgy of these steels. Much of this work has been summarized in review papers (Ref 1, 2, and 3) and will not be recapitulated here. A brief description of the metallurgical characteristics of these steels is necessary, however, for an understanding of their behavior. Maraging steels can be considered highly alloyed low-carbon, iron-nickel lath martensites. These alloys also contain small but significant amounts of titanium. Typical nominal compositions of maraging steels are given in Table 1 . The phase transformations in these steels can be explained with the help of the two phase diagrams shown in Fig. 1 , which depict the iron-rich end of the Fe-Ni binary system. Figure 1 (a) is the metastable diagram plotting the austenite-to-martensite transformation upon cooling and the martensite-to-austenite reversion upon heating. Figure 1 (b) is the equilibrium diagram showing that at higher nickel contents the equilibrium phases at low temperatures are austenite and ferrite. Table 1 Nominal compositions of commercial maraging steels Composition, %(a) Grade

Ni

Mo

Co

Ti

Al

Nb

Standard grades 18Ni(200)

18

3.3

8.5

0.2

0.1

...

18Ni(250)

18

5.0

8.5

0.4

0.1

...

18Ni(300)

18

5.0

9.0

0.7

0.1

...

18Ni(350)

18

4.2(b)

12.5

1.6

0.1

...

18Ni(Cast)

17

4.6

10.0

0.3

0.1

...

12-5-3(180)(c)

12

3

...

0.2

0.3

...

3.0

...

0.7

0.1

...

Cobalt-free and low-cobalt bearing grades Cobalt-free 18Ni(200)

18.5

Cobalt-free 18Ni(250)

18.5

3.0

...

1.4

0.1

...

Low-cobalt 18Ni(250)

18.5

2.6

2.0

1.2

0.1

0.1


Page 1225


Maraging Steels

01 Sep 2005

Cobalt-free 18Ni(300) 18.5 4.0 ... 1.85 0.1 ... (a) All grades contain no more than 0.03% C. (b) Some producers use a combination of 4.8% Mo and 1.4% Ti, nominal. (c) Contains 5% Cr

Fig. 1 Phase relationships in the iron-nickel system. (a) Metastable. (b) Equilibrium. Source: Ref 4

The metastable diagram indicates the typical behavior of these steels during cooling from the austenitizing or solution annealing temperature. No phase transformations occur until the Ms temperature, the temperature at which martensite starts to transform from austenite, is reached. Even very low cooling of heavy sections produces a fully martensitic structure, so there is no lack of hardenability in these alloys. Alloying elements alter the Ms temperature significantly, but do not alter the characteristic that transformation is independent of cooling rate. In addition to nickel, the other alloy elements present in maraging steels generally lower the martensite transformation range, with the exception of cobalt, which raises it. One of the roles of cobalt in maraging steels is to raise the Ms temperature so that greater amounts of other alloying elements (that is, titanium and molybdenum, which lower the Ms temperature) can be added while still allowing complete transformation to martensite before the steel cools to room temperature. Most grades of maraging steel have Ms temperatures of the order of 200 to 300 °C (390 to 570 °F) and are fully martensitic at room temperature. Therefore, retained austenite is generally not a problem in these alloys, and as a result, refrigeration treatments are not needed prior to aging. The martensite is normally a low-carbon, body-centered cubic (bcc) lath martensite containing a high dislocation density but no twinning. This martensite is relatively soft (~ 30 HRC), ductile, and machinable. The age hardening of maraging steels is produced by heat treating for 3 to 9 h at temperatures of the order of 455 to 510 °C (850 to 950 °F). The metallurgical reactions that take place during such treatment can be explained by using the equilibrium diagram (Fig. 1 b). With prolonged aging, the structure tends to revert to the equilibrium phases⎯primarily ferrite and austenite. Fortunately, the kinetics of the precipitation reactions that cause hardening are such that considerable age hardening⎯that is, approximately 20 HRC points (1035 MPa, or 150 ksi)⎯occurs before the onset of the reversion reactions that produce austenite and ferrite. With long aging times or high temperatures, however, hardness will reach a maximum and then will start to drop, as shown by the data in Fig. 2 . Softening in these steels usually results not only from overaging in the usual sense of the term⎯that is, coarsening of the precipitate particles⎯but also from austenite reversion. The two processes are interlinked; dissolution of metastable nickel-rich precipitate particles in favor of equilibrium iron-rich precipitates locally enriches the matrix in nickel, which favors austenite formation. Very substantial amounts of austenite (of the order of 50%) can eventually be formed by overaging. Fig. 2 Hardness of 18Ni(250) maraging steel versus aging time for various aging temperatures. Source: Ref 4


Page 1226


Maraging Steels

01 Sep 2005

Maraging steels normally contain little or no austenite after standard maraging heat treatments. However, austenite is always present in the heat-affected zones around welds and is sometimes deliberately formed to enhance fabricability or service performance. For example, if maraging steel is to be used in an application where overaging in service is expected, such as in dies for aluminum die castings, it is deliberately overaged slightly before being put into service. This minimizes overaging in service, which often produces tensile stresses at the surface. Extreme overaging to form large amounts of austenite has also been employed as an intermediate treatment to enhance response to cold working or to minimize the effects of thermal gradients during hot working and subsequent storage of extraordinary heavy sections. Age hardening in maraging steels results primarily from the precipitation of intermetallic compounds. Precipitation takes place preferentially on dislocations and within the lath martensite to produce a fine uniform distribution of coherent particles. The major hardener is molybdenum, which upon aging initially forms Ni3Mo, with an orthorhombic Cu3Ti-type structure (Ref 5). The metastable Ni3Mo phase forms initially because of its better lattice fit with the bcc martensitic matrix. Growth of the Ni3Mo is restricted by coherency strains, and as such, further aging results in the in situ transformation of Ni3Mo to the equilibrium Fe2Mo phase, which has a hexagonal C14-type structure (Ref 6). Titanium, which is generally present in maraging steels, promotes additional age hardening through the precipitation of Ni 3Ti, which has a DO24 ordered hexagonal structure (Ref 6). Cobalt does not directly participate in the age-hardening reaction, because this element does not form a precipitate with iron, nickel, molybdenum, or titanium in the 18Ni maraging alloy system. The main contribution of cobalt is to lower the solubility of molybdenum in the martensitic matrix and thus increase the amount of Ni3Mo precipitate formed during age hardening. Some hardening also results from a short-range ordering reaction in the matrix that involves cobalt. Molybdenum also plays the necessary supplemental role of minimizing localized grain-boundary precipitation by lowering the diffusion coefficients of a number of elements in solid solution. Precipitation of these grain-boundary phases severely impairs the toughness of most molybdenum-free ferrous alloys. Work by Schmidt (Ref 7) has shown that discrete particles of austenite are also present on the grain and subgrain boundaries in molybdenum-free 18Ni(300). It has been theorized that the precipitation of these discrete particles of austenite at the grain and subgrain boundaries results in a nickel-depleted zone, which adversely affects the toughness and ductility of the molybdenum-free 18Ni(300) alloy on a localized scale. The precipitate particles are of a lattice size comparable to that of the martensite matrix and cause little distortion of the matrix. This characteristic, together with the absence of carbon, allows the steel to be age hardened to very high strength levels while minimizing changes in the shape of the part being hardened. Another important precipitation reaction, which must be avoided if optimum toughness and ductility are to be achieved, can take place when the steel is in the austenitic condition. Maraging steels containing titanium are susceptible to the formation of TiC films at austenite grain boundaries. The solutioning of carbides at temperatures greater than 1150 °C (2100 °F), followed by holding at temperatures of the order of 900 to 1095 °C (1650 to 2000 °F), results in the reprecipitation of titanium carbide in the form of films on austenite grain boundaries. As discussed in the section "Hot Working" in this article, this problem first became evident in hot-worked billets and can be avoided by proper hot-working procedures. Prolonged annealing in this range should also be avoided.

Commercial Alloys Table 1 lists the chemical compositions of the more common grades of maraging steel. The nomenclature that has become established for these steels is nominal yield strength (ksi units) in parentheses. Thus, for example, 18Ni(200) steel is normally age hardened to a yield strength of 1380 MPa (200 ksi). The first three steels in Table 1 ⎯18Ni(200), 18Ni(250), and 18Ni(300)⎯are the most widely used and most commonly available grades. The 18Ni(350) grade is an ultrahigh-strength variety made in limited quantities for special applications. Two 18Ni(350) compositions have been produced (see the footnote in Table 1 ). The 18Ni(Cast) grade was developed specifically as a cast composition. Special varieties of maraging steels have been developed, including stainless grades, other cast grades, grades of other strength


Page 1227


Maraging Steels

01 Sep 2005

levels, a cobalt-free variety for nuclear service, a grade especially suited for heavy sections, and a grade with superior magnetic characteristics. Some of these steels have been made and used commercially, but only in limited amounts for specific applications. A number of cobalt-free maraging steels and a low-cobalt bearing maraging steel have recently been developed. The driving force for the development of these particular alloys was the cobalt shortage and resultant price escalation of cobalt during the late 1970s and early 1980s. The nominal compositions for these alloys are also listed in Table 1 . The wrought steels in Table 1 are produced in various forms: forgings, plate, sheet, and bar stock are generally available. Most applications use bars or forgings. Thin strip, although it has been produced, is commercially uncommon. Special varieties can be produced in shapes suitable for specific applications.

Processing Melting. Most grades of maraging steel are either air melted, then vacuum arc remelted, or vacuum induction melted, then vacuum arc remelted. Premium grades of maraging steels used in critical aircraft and aerospace applications, for which minimum residual element (carbon, manganese, sulfur, and phosphorus) and gas (O2, N2, and H2) contents are required, are triple melted using the air, vacuum induction, and vacuum arc remelting processes. The primary goals in melting are a clean microstructure, homogeneity, and low levels of deleterious residual elements, particularly sulfur and carbon. The sensitivity of maraging steels to the presence of inclusions is comparable to that of conventional steels of similar strength. Their sensitivity to residual elements, however, is different. Maraging steels show very little sensitivity to elements, such as arsenic, antimony, and phosphorus, that cause temper embrittlement in conventional quenched and tempered steels. As discussed in the section "Surface Treatment" in this article, maraging steels appear to be more resistant to hydrogen embrittlement than low-alloy steels. Carbon and sulfur are the most deleterious impurities in maraging steels because they tend to form brittle carbide, sulfide, carbonitride, and carbosulfide inclusions. These particles crack when the metal is strained, thus initiating fracture and lowering toughness and ductility. Steels with high titanium and molybdenum contents are prone to microsegregation of these elements during solidification. If the ingot structure is not homogenized during hot working, a banded microstructure with highly anisotropic mechanical properties may persist. The only effective remedy for microsegregation is to adjust the ingot size and hot-working schedule to ensure elimination of the segregation. With the 18Ni(350) grade, ingot size may have to be smaller than normal, and the maximum hot-working temperature should be kept below 1230 °C (2250 °F). Hot Working. Maraging steels can be hot worked by conventional steel mill techniques, even though allowances must be made for several unique characteristics. Steels with high titanium contents have greater hot strength than conventional steels and require higher hot-working loads or higher working temperatures. Working above about 1260 °C (2300 °F) should be avoided. To maximize their mechanical properties, maraging steels should be hot worked at the lowest temperatures that equipment power limitations permit. The precipitation of TiC films at austenite grain boundaries must be avoided. This phenomenon first came to light in billets that had been worked at very high temperatures and then allowed to cool slowly through the temperature range of 750 to 1095 °C (1380 to 2000 °F) or to cool with inadvertent thermal arrests in this range. It is essential that long dwell times in this temperature range be avoided after working is completed so that the titanium and carbon remain in solution. However, one should keep in mind that it is safe to heat into the 750 to 1095 °C (1380 to 2000 °F) range from room temperature because stable carbides will already have precipitated. This temperature range should be avoided only when cooling from temperatures above 1150 °C (2100 °F). Cold Working. Maraging steels can be cold worked by any conventional technique when in the solution-annealed (unaged or as-transformed) condition. They have very low work-hardening rates and can be subjected to very heavy reductions (>50%) with only slight accompanying gains in hardness. In maraging steels, both the strain-hardening exponent and tensile elongation are relatively low. Thus, cold-forming operations that depend on uniform tensile strain before necking (such as deep drawing) are generally limited to mild degrees of deformation. Uniform elongation can be significantly improved by overaging, which produces substantial amounts of austenite. This austenite is metastable and transforms to martensite during cold working, which greatly increases uniform elongation and deep-drawing capabilities. Heavily cold worked structures can be softened by austenitizing or solution annealing at temperatures of about 815 °C (1500 °F). Cold-worked pieces can be directly age hardened, in which case the total strength includes the increase in strength produced by cold working plus that produced by precipitation hardening. Large amounts of prior cold work decrease toughness after direct aging, and considerable anisotropy in elastic modulus and in toughness may be present in unidirectionally worked structures. Machining. Maraging steels can be machined by any conventional technique when in the solution-annealed or age-hardened condition. Machinability is generally as good as or better than that of conventional steels of the same hardness. Rigid equipment and firm tool supports are necessary. Tools should be flooded with cutting fluids free of lead, sulfur, and other low-melting components, because residues of these elements may cause surface embrittlement during subsequent heat treatment. The new generation of ceramic tooling provides superior cutting speeds and tool life when working with 18Ni maraging steels. Heat Treating. The properties of conventionally heat treated maraging steels are given in Table 2 and the mechanical properties of 18Ni(250), cobalt-free 18Ni(250), and low-cobalt bearing 18Ni(250) are compared in Table 3 . Maraging steels are usually solution annealed (austenitized) 1 h for each 25 mm (1 in.) of section size. It may be necessary to heat treat sheet products


Page 1228


Maraging Steels

01 Sep 2005

in dry hydrogen or dissociated ammonia to minimize surface damage. The cooling rate after annealing has a minimal effect on microstructure or properties. It is essential, however, to cool the steel completely to room temperature before aging. If this is not done, the steel may contain untransformed austenite and may be much softer than expected. Table 2 Heat treatments and typical mechanical properties of standard 18Ni maraging steels Tensile strength

Yield strength ksi


Reduction in area, %

1400

203

10

60

155−240

140−220

260

1700

247

8

55

120

110

297

2000

290

7

40

80

73

355

2400

348

6

25

35−50

32−45

Heat treatment(a )

MPa

ksi

MPa

18Ni(200)

A

1500

218

18Ni(250)

A

1800

18Ni(300)

A

2050

18Ni(350)

B

2450

Grade

Fracture toughness p p MPa m ksi in:

18Ni(Cast) C 1750 255 1650 240 8 35 105 95 (a) Treatment A; solution treat 1 h at 820 °C (1500 °F), then age 3 h at 480 °C (900 °F). Treatment B: solution treat 1 h at 820 °C (1500 °F), then age 12 h at 480 °C (900 °F). Treatment C: anneal 1 h at 1150 °C (2100 °F), age 1 h at 595 °C (1100 °F), solution treat 1 h at 820 °C (1500 °F) and age 3 h at 480 °C (900 °F)

Table 3 Comparison of the longitudinal, room-temperature mechanical properties of standard, cobalt-free, and low cobalt-bearing 18Ni(250) maraging steels Heat treatment: solution heat 1 h at 815 °C (1500 °F), then age 5 h at 480 °C (900 °F). Testing was conducted on 63.5 × 88.9 mm (2.5 × 3.5 in.) billets produced from 200 mm (8 in.) diam vacuum induction melted/vacuum arc remelted ingots. Charpy V-notch Ultimate 0.2% offset Reductio impact Plane-strain fracture tensile yield Elongation n toughness(a) toughness(a) strength strength in 25 mm in area, p p MPa m (1 in.), % % Grade MPa ksi MPa ksi J ft · lbf ksi in: 18Ni(250)

1870

271

1825

265

12

64.5

37

27

138

125

Cobalt-free 18Ni(250)

1895

275

1825

265

11.5

58.5

34

25

127

115

11

63.5

43

32

149

135

Low-cobalt 18Ni(250) 1835 266 1780 258 (a) Longitudinal-short transverse orientation tested (L-S orientation)

Aging is normally done at 480 °C (900 °F) for 3 to 6 h. For certain applications (such as die casting dies), aging at about 530 °C (985 °F) is employed. Figure 2 shows hardness versus aging time for 18Ni(250). Hardening is initially very rapid; several minutes at temperature will cause a substantial boost in strength in all grades of maraging steel. With very long times at temperature, hardness begins to drop due to coarsening of the precipitate particles and reversion of the martensite matrix to austenite. Austenite generally starts to form as rather small particles at both prior austenite boundaries and lath-martensite boundaries. Thermal cycling of maraging steels between the Mf temperature (temperature at which martensite formation finishes during cooling) and a temperature considerably in excess of the solution-heating temperature can be used to refine the grain structure of coarse-grain maraging steels. For example, Saul, Roberson, and Adair (Ref 8) were able to refine an ASTM grain size of −1/1 to an ASTM grain size of 6/7 in 18Ni(300) following three thermal cycles between room temperature and 1025 °C (1880 °F). The shear strains produced by the diffusionless shear transformations of martensite to austenite and of austenite to martensite provide the driving force for recrystallization during these thermal cycles. One should recall that grain sizes finer than ASTM 6/7 cannot be achieved by this process and that the process becomes less effective as the starting grain size becomes finer. The standard aging treatments listed in Table 2 produce contraction in length of 0.04% in 18Ni(200), 0.06% in 18Ni(250), and 0.08% in both 18Ni(300) and 18Ni(350). These very small dimensional changes during aging allow many maraging steel components to be finish machined in the annealed condition, then hardened. When precise dimensions must be held, an allowance for contraction can be made. Much effort has been devoted to examining the properties of overaged maraging steels. The general belief is that a microstructure containing coarse precipitate particles and finely distributed austenite particles should have good resistance to both fracture and stress-corrosion cracking. In many instances, this has been found to be so, and very impressive improvements in plane-strain fracture toughness, KIc, or threshold stress intensity for stress-corrosion cracking, KIscc, have been achieved with only modest reductions in yield strength. Unfortunately, however, there appears to be more heat-to-heat variability when overaging heat treatments are used. This variability is the result of a greater sensitivity of austenite formation to minor changes in composition and processing. Even the most homogeneous-looking structures contain minor alloy segregation that is manifested in the form of alloy-rich and alloy-lean bands. The alloy-rich bands tend to overage more rapidly than their alloy-lean counterparts. It is therefore difficult to recommend specific overaging heat treatments that will produce consistent mechanical properties. Generally, if a specific yield strength is required, it is better to use a maraging steel in which the required strength can be produced by conventional aging than to use an overaged steel of higher strength.


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Maraging Steels

01 Sep 2005

Surface Treatment. Grit blasting is the most efficient technique for removing oxide films formed by heat treatment. Maraging steels can be chemically cleaned by pickling in sulfuric acid or by duplex pickling in hydrochloric acid and then in nitric acid plus hydrofluoric acid. As with conventional steels, care must be taken to avoid over-pickling. The sodium hydride cleaning of maraging steels should be avoided to minimize problems with crack formation. Grease and oils can be removed by cleaning in trichloroethane-type solutions. Maraging steels can be nickel plated in chloride baths provided that proper surface-activation steps are followed. Heavy chromium deposits can be plated on top of nickel electrodeposits. Maraging steels are less susceptible to hydrogen embrittlement during plating than conventional quenched and tempered steels of comparable hardness. They are not immune to hydrogen, however, and baking after plating is recommended. Baking should be done at temperatures of about 150 to 205 °C (300 to 400 °F) for periods of 3 to 10 h, depending on size and baking temperature. Baking cannot be combined with age hardening, because considerable hydrogen remains in the steel after heat treating at the higher temperatures. Considerable surface hardening can be achieved by nitriding maraging steels in dissociated ammonia. Hardness levels equivalent to 65 to 70 HRC can be achieved at depths of up to 0.15 mm (0.006 in.) after nitriding for 24 to 48 h at 455 °C (850 °F). Nitriding at this temperature allows age hardening to occur during nitriding; therefore, the two processes can be accomplished simultaneously. Salt bath nitriding for 90 min at 540 °C (1000 °F) has been done successfully. Such treatment must be very carefully controlled to avoid excessive overaging. Both the fatigue strength and the wear resistance of maraging steels are improved by nitriding. Welding. One of the main virtues of maraging steels is their excellent weldability. The extensive research and development work that has been done to optimize weldability is reviewed in Ref 4 and 9; only a brief summary of that work is given here. Perhaps the characteristic that most enhances the weldability of maraging steels is that their low carbon content produces a soft, ductile martensite on cooling. A weld heat-affected zone in maraging steels can be divided into three regions. The region closest to the fusion line contains coarse martensite produced by solution annealing. Next is a narrow region containing reverted austenite produced by heating into the 595 to 805 °C (1100 to 1480 °F) range. Finally, there is a region where the maximum temperature reached during welding ranges from ambient temperature up to 595 °C (1100 °F); this region contains martensite that has been age hardened by various amounts. As shown in Fig. 3 , the heat-affected zone in an as-welded structure is relatively soft. Because the metal in the area immediately surrounding the weld is soft and ductile, residual stresses are low, and the susceptibility to weld cracking is considerably less than in steels hardened by quenching. Subsequent local aging brings the hardness of the weld zone up to that of the base metal; this effect is also shown in Fig. 3 . Resolution heat treating is advisable after welding if optimum properties are desired in the heat-affected weld zone. Fig. 3 Microhardness of a weld heat-affected zone in 18Ni(250) maraging steel. Source: Ref 4

All conventional welding processes have been used for welding maraging steel. Brazing can also be done. Most fabrication thus far, however, has employed inert-gas welding processes; the least work has been done with flux-shielded processes (Ref 4, 9). Typical compositions of filler wire for welding maraging steels are listed in Table 4 . These compositions are very similar to those of the steels with which they are used. After welding, an aging treatment is used to increase the strength of the weld joint. Weld metal strength does not depend significantly on the welding process used to make the joint; in most applications, joint efficiencies exceeding 90% can be achieved. Table 4 Typical compositions of filler wire for welding 18Ni maraging steels Composition of filler wire, %(a)

Base metal grade

Ni

Mo

Co

Ti

Al

18Ni(200)

18

3.5

8

0.25

0.10

18Ni(250)

18

4.5

8

0.50

0.10

18Ni(300)

18

4.5

10

0.80

0.10


Page 1230


18Ni(350) (a) All grades contain no more than 0.03% C.

Maraging Steels

17

3.7

01 Sep 2005

12.5

1.60

0.15

The toughness of the heat-affected zone after age hardening usually matches that of the unaffected base metal. The toughness of the weld metal, however, depends on the joining process. Gas tungsten arc welding produces the best toughness, gas metal arc welds have somewhat lower toughness, and the toughness of flux-shielded welds is poorer yet. In general, high-energy processes that produce coarser weld beads will produce weld metal of lower toughness. This is of special concern with submerged arc welding, where low toughness negates the advantage of the high welding speed associated with that process. The general rules that follow are helpful for obtaining the best properties: • • • • • •

Avoid prolonged dwell times at high temperatures Avoid preheat Keep interpass temperatures below 120 °C (250 °F) Minimize weld energy Avoid slow cooling rates Keep welds as clean as possible

Powder Metallurgy Products. Maraging steels having yield strengths of about 1585 MPa (230 ksi), combined with reasonable ductility, have been made by sintering elemental powders. Gas-atomized alloy powders that have been canned and extruded show excellent ductility at yield strengths up to 1900 MPa (275 ksi). However, during the atomization and collection of atomized maraging steel powder, care must be taken to minimize contamination of the surface of the powder particles. The presence of oxides or other nonmetallic inclusions on these surfaces can adversely affect the mechanical properties of parts consolidated from contaminated powder. Powder metallurgy parts made of maraging steel have been commercially produced only in limited quantities.

Properties of Maraging Steels Mechanical Properties. Typical tensile properties and fracture toughness values of the conventional grades of maraging steel are listed in Table 2 . One of the distinguishing features of maraging steels is superior toughness compared to conventional steels. Figure 4 compares KIc values for several maraging steels with those of quenched and tempered alloy steels as a function of tensile strength. The toughness of maraging steels is sensitive to purity level, and carbon and sulfur levels in particular should be kept low to obtain optimum fracture toughness. Fig. 4 Plane-strain fracture toughness of maraging steels compared with fracture toughness of several ultrahigh strength steels as a function of tensile strength. Source: Ref 2

The mechanical properties of 18Ni(250), cobalt-free 18Ni(250), and low-cobalt bearing 18Ni(250) are compared in Table 3 . These data show that all three alloys exhibit comparable strength when aged at 480 °C (900 °F) for 5 h. However, these data also show that, in general, the cobalt-free 18Ni(250) material displays reduced localized ductility as measured by the percentage of reduction in area, and reduced plane-strain fracture toughness compared to the standard 18Ni(250) and low-cobalt bearing


Page 1231


Maraging Steels

01 Sep 2005

18Ni(250) alloys. The reduced ductility and toughness of the cobalt-free 18Ni(250) alloy is related to the absence of cobalt in this material. The lack of cobalt eliminates the previously discussed cobalt/molybdenum interaction and thus necessitates a higher level of titanium, which is a more potent embrittling agent than either cobalt or molybdenum when present at levels significantly in excess of 1.3%. The effects of temperature on the mechanical properties of maraging steels are illustrated in Fig. 5 . Maraging steels can be used for prolonged service at temperatures up to approximately 400 °C (750 °F). Yield and tensile strengths at 400 °C (750 °F) are about 80% of the room-temperature values. Long-time stress-rupture failures can occur at 400 °C (750 °F), but the rupture stresses are fairly high. At temperatures above 400 °C (750 °F), reversion of the martensite matrix to austenite becomes dominant, and long-term load-carrying capacities decay fairly quickly. Fig. 5 Effect of temperature on the mechanical properties of 18Ni maraging steels. (a) Stress. (b) Ductility. (c) Hardness. Source: Ref 10


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Maraging Steels

01 Sep 2005

At cryogenic temperatures, the strengths of maraging steels increase similarly to those of steels (Fig. 5 ); KIc values fall off


Page 1233


Maraging Steels

01 Sep 2005

roughly in a linear fashion with decreasing temperatures. The fatigue properties of maraging steels are comparable to those of other high-strength steels. Smooth-bar and notched-bar fatigue properties for 18Ni(200), 18Ni(250), and 18Ni(300) grades are summarized in Fig. 6 . Fatigue crack growth rates in maraging steels obey the da/dN = (∆K)m relationship commonly observed in steels, and the rates are similar to those of conventional steels. The bulk of the fatigue life in maraging steels is in the crack initiation stage. In general, cracks tend to initiate at noncoherent intermetallic inclusions, which once again suggests that low levels of residual elements and clean melt practices are important in this alloy system. Improved fatigue properties can be obtained by shot peening and by nitriding. Fig. 6 Rotating-bean fatigue properties of three 18Ni maraging steels. Source: Ref 2

Resistance to Corrosion and Stress Corrosion. Maraging steels have slightly better corrosion resistance than tempered martensitic alloy steels. In industrial and marine atmospheres, the corrosion rates of maraging steels are about half those of conventional steels. In static and flowing seawater, maraging and conventional steels have essentially the same corrosion rates. In saline and acidic solutions, maraging steels show somewhat better corrosion resistance. Their hot oxidation resistance is noticeably better than that of tempered martensitic alloy steels because of the nickel and cobalt contents of these materials. In general, contact with more noble metals should be avoided to minimize the likelihood of galvanic corrosion. Maraging steels are susceptible to stress-corrosion cracking in most aqueous environments; their resistance to cracking increases significantly at lower yield strengths. Maraging steels have better resistance to stress-corrosion cracking than tempered martensitic steels of comparable strength. Figure 7 compares KIscc values of several high-strength steels. The resistance of maraging steels to stress-corrosion cracking generally parallels fracture toughness. Processing techniques that improve KIc (such as vacuum melting, proper hot working, and keeping residual impurity levels low) also improve resistance to stress-corrosion cracking. Surface treatments such as painting or shot peening have been successfully used to increase resistance. Cathodic protection can also be used, but must be employed very carefully to avoid hydrogen charging that could initiate cracking. Maraging steels can be used in aqueous environments; however, as with all high-strength steels, proper precautions to avoid stress-corrosion cracking must be taken. Fig. 7 Threshold stress intensity for stress-corrosion cracking (KIscc) values for maraging steels and other high-strength steels as a function of yield strength. Source: Ref 11


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Maraging Steels

01 Sep 2005

A comparison of the environmentally assisted cracking (EAC) behavior of production scale lots of 18Ni(250), low-cobalt 18Ni(250), and cobalt-free 18Ni(250) is shown in Fig. 8 , 9 , and 10 . The 18Ni(250) and low-cobalt 18Ni(250) materials used for this comparison testing were produced by Carpenter Technology Corporation; the cobalt-free 18Ni(250) was produced by Teledyne Vasco Corporation. The KIc and KIscc data contained in Fig. 8 and 9 were produced by tests using compact tension specimens. The KIscc testing was conducted on samples mounted in a fiberglass rack that was 45° to the vertical and facing the p prevailing southeast wind on a floating pier located in Galveston Bay. All samples were loaded to Kinit of 66 MPa m (60 ksi p in:), and the test duration was 1000 h. The proof-ring tensile data contained in Fig. 10 is from a test conducted in a stagnant 3.5% NaCl solution using Plexiglas vessels. Figures 8 and 9 show that both the 18Ni(250) and low-cobalt 18Ni(250) alloys display improved KIc, KIscc, and EAC growth compared to the cobalt-free 18Ni(250) alloy. Based on Fig. 10 , the apparent threshold stress levels for 18Ni(250), low-cobalt 18Ni(250), and cobalt-free 18Ni(250) in stagnant 3.5% NaCl are 1035 MPa (150 ksi), 860 MPa (125 ksi), and 515 MPa (75 ksi), respectively. Fig. 8 Bar graphs comparing plane-strain fracture toughness KIc (a) and KIscc (circumferential-radial specimen orientation) (b) of low-cobalt bearing, standard, and cobalt-free 18Ni(250). KIscc testing was conducted in a marine atmosphere. Heat treatment: 815 °C (1500 °F), 1 h, air-cooled; 480 °C (900 °F), 5 h, air-cooled. Source: Ref 12

Fig. 9 Bar graph showing the variation in total crack length between low-cobalt bearing, standard, and cobalt-free 18Ni(250) maraging steels following KIscc testing in a marine atmosphere. Heat treatment: 815 °C (1500 °F), 1 h, air-cooled; 480 °C (900 °F), 5 h, air-cooled. Source: Ref 12


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Maraging Steels

01 Sep 2005

Fig. 10 Plot of stress level versus time to failure to compare behavior of low-cobalt bearing, standard, and cobalt-free 18Ni(250) maraging steels tested in stagnant 3.5% NaCl for 1000 h using proof-ring tensile specimens. Heat treatment: 815 °C (1500 °F), 1 h, air-cooled; 480 °C (900 °F), 5 h, air-cooled. Source: Ref 12

Physical Properties. Table 5 summarizes the physical properties of 18Ni(250) maraging steel. Grade 18Ni(250) is the most widely used and most thoroughly evaluated maraging steel. Consequently, more extensive physical property data are available for this steel than for other maraging steels (Ref 13). Table 5 Selected physical and mechanical properties of 18Ni(250) maraging steel Density, g/cm3 (lb/in.3)

8.0 (0.289)

Coefficient of thermal expansion, µm/m · K (µin./in. · °F) at 24−284 °C (75−543 °F)

10.1 (5.6)

Thermal conductivity, W/m · K (Btu · in./ft3 · h · °F) Ã20 °C (70 °F)

19.7 (137)

Ã50 °C (120 °F)

20.1 (139)

Ã100 °C (212 °F)

20.9 (145)

Electrical resistivity, µΩ · m ÃAnnealed

0.6−0.7

ÃAged

0.36−0.6


Page 1236


Maraging Steels

01 Sep 2005

Melting temperature, °C ( °F)

1430−1450 (2605−2640)

Length change during aging, %

−0.06

Poisson's ratio

0.30

Elastic modulus, GPa (106 psi)

186 (27)

Shear modulus, GPa (106 psi)

71 (10.3)

Magnetic coercive force, A · m−1 (Oe) ÃAnnealed

1750−2700 (22−34)

ÃAged

1670−4300 (21−54)

Applications Maraging steels have been used in a wide variety of applications, including missile cases, aircraft forgings, structural parts, cannon recoil springs, Belleville springs, bearings, transmission shafts, fan shafts in commercial jet engines, couplings, hydraulic hoses, bolts, punches, and dies. Maraging steels have been extensively used in two general types of applications: • Aircraft and aerospace applications, in which the superior mechanical properties and weldability of maraging steels are the most important characteristics • Tooling applications, in which the excellent mechanical properties and superior fabricability (in particular, the lack of distortion during age hardening) are important In many applications, even though maraging steels are more expensive than conventional steels in terms of alloy cost, finished parts made of maraging steel are less expensive because of significantly lower fabrication costs. Therefore, it is often economics rather than properties alone that determine the use of maraging steels. REFERENCES 1. S. Floreen, Metall. Rev., Vol 13 (No. 126), 1968 2. A. Magnee, J.M. Drapier, J. Dumont, D. Coutsouradis, and L. Habraken, "Cobalt-Containing High-Strength Steels," Cobalt Information Center, 1974 3. A.M. Hall and C.J. Slunder, "The Metallurgy, Behavior, and Application of the 18 Percent Nickel Maraging Steels," Report SP5051, National Aeronautics and Space Administration, 1968 4. F.H. Lang and N. Kenyon, Bulletin 159, Welding Research Council, 1971 5. W.B. Pearson, Handbook of Lattice Spacings and Structure of Metals and Alloys, Vol 2, Pergamon Press, 1967 6. W.B. Pearson, Handbook of Lattice Spacings and Structure of Metals and Alloys, Vol 1, Pergamon Press, 1958 7. M.L. Schmidt, Maraging Steels: Recent Developments and Applications, The Minerals, Metals & Materials Society, 1988, p 213−235 8. G. Saul, J.A. Roberson, and A.M. Adair, in Source Book on Maraging Steels, American Society for Metals, 1979 9. C.R. Weymueller, How to Weld Maraging Steels, Weld. Des. Fabr., Feb 1989 10. "18% Ni Maraging Steel Data Bulletin," The International Nickel Company, Inc., 1964 11. D.P. Dautovich and S. Floreen, in Proceedings of the International Conference on Stress Corrosion Cracking and Hydrogen Embrittlement, National Association of Corrosion Engineers, 1976 12. M.L. Schmidt, Carpenter Technology Corporation, unpublished research 13. R.P. Wei, Ferrous Alloys, in Aerospace Structural Metal Handbook, Belfour Stulen, Inc., 1970


Page 1237


Ferrous Powder Metallurgy Materials

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Specialty Steels and Heat-Resistant Alloys Ferrous Powder Metallurgy Materials Revised by Leander F. Pease III, Powder-Tech Associates, Inc. POWDER METALLURGY (P/M) in its simplest form consists of compressing metal powders in a shaped die to produce green compacts. These are then sintered, or diffusion bonded, at elevated temperatures in a furnace with a protective atmosphere. During sintering, the constituents usually do not melt, and the compacts become substantially strengthened by the development of bonds between individual particles. For a specific metal powder and sintering condition, increased compact density results in improved mechanical properties. The density of sintered compacts may be increased by re-pressing. When re-pressing is performed primarily to increase dimensional accuracy rather than density, it is termed sizing. When re-pressing is intended to change the contour of the surface in contact with the punches, it is termed coining. For example, a sintered blank could be coined so that the surface is indented with small slots or letters and numbers. The re-pressing may be followed by resintering, which relieves the stresses due to cold work and may further strengthen the compact. Alloy compacts can be formed from mixtures of metal powders that completely or partially diffuse during sintering. Alternatively, each individual particle may be completely prealloyed prior to compaction. The diffusion bonding process may be accelerated by sintering at a temperature at which one of the constituent metals is molten. For metals with relatively high melting points, such as iron or tungsten, a skeleton may be pressed, which is then infiltrated by a molten metal, such as copper or silver, having a melting point lower than that of the skeleton. Process Capabilities. Theoretical density commonly refers to the density of a pore-free material of defined chemistry and thermomechanical history. For example, the theoretical density of iron or low-carbon steel is about 7.87 g/cm3. Adding carbon or oxide inclusions to the solid steel lowers the theoretical density, even though no pores are formed. By pressing and sintering only, parts are produced at 80 to 93% of theoretical density. By re-pressing, with or without resintering, the materials may be further densified to 85 to 96% of theoretical density. High-temperature sintering will also produce parts at these high densities. The density of pressed parts is limited by the size and shape of the compact. The most common P/M materials for structural parts are iron-copper-carbon, iron-nickel-carbon, and iron-carbon. Parts made from these materials respond to heat treatment with a defined hardenability band. Iron parts that are low in carbon and high in density can be carburized and quenched to form a definite, hard case. Powder metallurgy parts are frequently competitive with forgings, castings, stampings, numerically controlled machined components, and fabricated assemblies. Within the limitations inherent to the P/M processes (discussed in detail in Powder Metallurgy, Volume 7 of the 9th Edition of Metals Handbook), parts can be fabricated to final or near-final shape, thereby eliminating or reducing scrap metal, secondary machining, and assembly operations. Although the unit cost of metal powder (about $0.73/kg, or $0.33/lb in 1989) is usually higher than that of steel bars, the savings achieved by eliminating fabricating operations and minimizing scrap losses often result in lower total costs for P/M parts. The pressing operation in an intricate set of tooling provides very complex geometry, at production rates of 8 to 60 pieces per minute. Certain metal products can be produced only by powder metallurgy; among these products are materials whose porosity (number, distribution, and size of pores) is controlled. Two examples of controlled-porosity materials are filter elements and self-lubricating bearings. The porosity in a self-lubricating bearing is filled with oil, which lubricates the bearing/shaft interface as the temperature starts to rise. Successful production by powder metallurgy depends on the proper selection and control of process variables: • • • • • • • •

Powder characteristics Powder preparation Type of compacting press Design of compacting tools and dies Type of sintering furnace Composition of the sintering atmosphere Choice of production cycle, including sintering time and temperature Secondary operations and heat treatment


Page 1238



01 Sep 2005

Metal Powder Characteristics and Control The size and shape of powder particles have important effects on pressing and sintering characteristics, but the performance of each type of powder can be determined only empirically. The size and shape of iron powder particles are largely determined by the process used in producing the powder. Powder particles representative of the principal production processes are shown in Fig. 1. Fig. 1 Scanning electron micrographs of various iron powder particles. (a) Water-atomized and annealed iron powder (Ancorsteel 1000). Arrows indicate small fines that were agglomerated into larger particles. 190× (b) Iron powder (Atomet 28). Arrows indicate porosity in the spongy regions. 750×. (c) Hydrogen-reduced sponge iron (Pyron 100). Arrows indicate pores opening into the spongy interior. 1000×. (d) Carbon-reduced iron ore (MH-100). Arrows indicate one particle with coarse internal porosity. 750×

To ensure uniformity in the handling and press performance of powders from batch to batch and in the strength of the compacts formed from them, laboratory tests are usually made on samples of powder from each batch. The results of quality control tests can be applied only to samples of the same basic composition made by the same method. This limitation is evident in Fig. 2 , which shows typical compressibility values for three types of iron powders, atomized, prealloyed 4600V, and hydrogen-reduced sponge iron. Fig. 2 Effects of compacting pressure on density and mechanical properties. (a) Density versus compacting pressure for three types of iron powders. (b) Ultimate tensile strength versus density for FC-0208 (Fe-2%Cu-0.8%C).


Page 1239



01 Sep 2005

Test procedures are commonly specified for sampling and for the determination of particle size distribution, flow rate, apparent density, weight loss in hydrogen, compressibility, green strength, sintering characteristics, and chemical composition. A cross-index of standards for testing metal powders and ferrous P/M products is given in Table 1 . Table 1 Test methods applicable to powder metallurgy Designation Methods applicable to metal powders

ASTM

MPIF

ISO

Compressibility of metal powders

B 331

...

...

Compression testing of metallic materials at room temperature

E9

...

...

Density, apparent, of metal powders

B 212

4

3923/1

Density, apparent, of nonfree-flowing metal powders

B 417

28

3923/2

Flow rate of metal powders

B 213

3

4490

Hydrogen loss of metal powders

E 159

2

449/3

Insoluble matter in iron and copper powders

E 194

6

4496

Iron content of iron powder Particle size average of metal powders by Fisher subsieve sizer

...

7

...

32

... ...

Sampling finished lots of metal powders

B 215

1

3954

Sieve analysis of granular metal powders

B 214

5

4497

Subsieve analysis of granular metal powders by air classification

B 293

12

Tension test specimens for pressed and sintered metal powders

E8

10

Recommended practice for analysis by microscopical methods for particle size distribution of particulate substances of subsieve sizes

E 20

...

Terminology

B 438

...

2739

Tap density of metal powders

B 527

...

3953


... 2740 ...

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01 Sep 2005

Lubricant content of mixed powders Compactability of metal powders

... B 331

...

4495

45

...

Apparent density of Arnold meter

...

48

...

Test of copper base powder for infiltrating

...

49

...

Designation Methods applicable to compacted and sintered materials Density and interconnected porosity

ASTM B 328

Determination of apparent hardness of sintered metal powder products Determination of dimensional change of sintered metal powder specimens Determination of transverse rupture or bending strength of sintered metal powder test specimens

42

2737, 2738

43

4498/1

44

4492

...

40

5754

41

3325

37

4507

B 528

Determining the case hardness of powder metallurgy parts

ISO

... B 610

Determination of impact strength of sintered metal powder specimens

MPIF

...

Green strength of compacted metal powder specimens

B 312

15

3995

Determination of radial crushing strength of bearings

B 438

...

2739

...

3312

...

3369

Determination of Young's modulus Determination of density of impermeable materials

... B 311

Fatigue test pieces

...

...

3928

Determination of bubble test pore size in filters

...

39

4003

Determination of fluid permeability

...

...

4022

Sampling. Metal powders in containers are usually sampled by turning a specially designed auger-type sampler called a sampling thief vertically to the bottom of the container and removing the column of powder. (A straight, hollow tube does not provide a representative sample). One or more of these samples are taken from each drum of powder, the number of samples depending on the drum size. Samples taken in this manner are combined and reduced to the required size by means of a so-called sample splitter, or by rolling, coning, and quartering. Particle Size Distribution. A complete screen analysis of the particle size distribution is determined by placing a 100 g (3.5 oz) sample of powder in the top screen of a "nest" of standard sieves having successively smaller openings, usually 80, 100, 140, 200, and 325 mesh. The assembly is shaken mechanically for 15 min. The powder that is retained on each mesh size and the powder that passes through the finest sieve are reported as weight percentages of the sample. The powder passing 325-mesh can be further analyzed for its subsieve particle size distribution by use of an optical or electron microscope or by the scattering of laser light. Flow Time. The standard (Hall) flowmeter for metal powders is a funnel-shape apparatus having a calibrated orifice. The time required for 50 g (1.75 oz) powder to flow unaided through the orifice is reported as the flow time of the powder. Apparent Density. The Hall flowmeter is also used to determine the apparent density of free-flowing powders. For nonflowing powders, use is made of a baffled rectangular tower fed by a funnel having a large orifice (Scott volumeter). The Arnold meter measures apparent density and approximates the value that is found in actual press operation. All three instruments serve to supply a controlled flow of powder into a cup of known volume. The content of the full cup is weighed, and the apparent density is reported as g/cm3. Oxide Content. A fairly accurate indication of the hydrogen-reducible oxide content of metal powders can be obtained by reducing a weighed sample of the powder in hydrogen under standard time-temperature conditions. The sample must be cooled to room temperature in an atmosphere of dry hydrogen. Loss in weight caused by reduction of the oxides present is expressed as a percentage of the initial sample weight. Direct determination of oxygen may also be used to estimate the oxide content of the powder. Compressibility is the density to which a powder can be pressed at a given pressure. Alternatively, it can be expressed as the pressure required to attain a given density (see Fig. 2 ). Green Strength. The strength of a green (unsintered) compact is usually determined by pressing a given weight of powder to a specified density in the shape of a rectangular bar of standard size. This test bar is supported at both ends as a simple beam and loaded transversely to fracture. The modulus of rupture is reported as the green strength of the powder. Sintering Characteristics. A weighed quantity of a powder is compacted under controlled conditions in a die of known dimensions. After sintering under controlled conditions, dimensional measurements are again taken. The difference in size is usually given as a percentage of the die dimension. Finally, the specimen is tested for mechanical properties. Sintering characteristics are generally reported in terms of dimensional changes, final density, and mechanical properties. The test is usually done by comparing the dimensional change or strength to a standard or reference bar processed at the same time. Chemical Composition. In addition to the above tests, routine chemical analyses usually are run on all powders to determine the amount of alloying elements and impurities that might adversely affect the final product.


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01 Sep 2005

Powder Preparation Powders are mixed or blended before use. This is done to obtain specific properties (for example, dimensional change or strength) in the finished product. Mixing or blending must be carefully controlled. Also, it is important to use equipment that mixes rapidly and produces a uniform distribution of the powders in the desired proportions in minimum time. Excessive blending, especially in an overloaded blender, may work harden the powder, making it less compressible and thus more difficult to compact. The cascading of powder inside the blender will gradually round the edges of the particles, changing their shape and raising apparent density. As a result, the powder will exhibit different molding characteristics. The double-cone blender is probably the most extensively used of the various types available. The final operation before compaction is to mix the powders with a lubricant, such as synthetic wax or stearic acid. The lubricant not only minimizes die friction and wear, but also reduces interparticle friction, allowing the particles to pack more closely and resulting in higher density for the lower range of molding pressure. Because of a hydraulic effect, densities above 7.0 g/cm3 require less lubricant. Small percentages of alloying additives such as graphite, copper, or nickel powders, if required, are also introduced during final mixing. Powder manufacturers are beginning to provide bonded mixes in which 1=8% of a polymer attaches the additives to the iron powder. This prevents segregation, increases flow rate, and decreases dimensional variation among parts.

Powder Compacting Compacting powder serves several important functions: The powder is consolidated into the desired shape⎯a compact that must be strong enough to withstand subsequent processing; compacting controls the amount and type of porosity of the finished product; and compacting is largely responsible for the final dimensions of the part, subject to dimensional changes during sintering. Of the numerous compacting methods, closed-die pressure compacting and room temperature is the method that is most commonly used in the production of ferrous powder metallurgy parts. A detailed discussion of the various compacting methods, their capabilities and limitations, and the equipment required by each method may be found in Powder Metallurgy, Volume 7 of the 9th Edition of Metals Handbook or in Ref 1 and 2. Green strength is an important property of a green compact; the compact must be strong enough to avoid damage during ejection from the die and during normal handling between compacting and sintering. The green strength of a compact is affected by the type and composition of the powder, the distribution of powder particles with respect to both size and shape, the quantity and type of lubricant added to the powder, the configuration of the die, and the compacting pressure. If other compacting conditions are held constant, the green strength of powder metallurgy compacts will vary directly with compacting pressure. Also, as shown in Fig. 2 , green density of compacts increases with compacting pressure. Green density is indicative of the density that can be expected after sintering.

Sintering Sintering is the process by which a compact of metal powder is bonded by heating at a temperature below the melting point, or liquidus, of the major constituent. In iron-graphite mixes, no melting occurs during sintering. In the commonly Fe-2Cu-0.8C mixes, the copper melts and diffuses into the iron. For an M-2 tool steel, sintering is done above the solidus temperature with 15 to 20% permanent liquid phase. Densification does not occur in most commercially produced P/M parts. Sintering conditions and alloys are adjusted so that the final part is very close to the die size that molded it. Sometimes a liquid phase forms and assists in sintering. This is true for the Fe-2Cu-0.8C mixtures noted above. The Fe-0.45P and Fe-0.8P contain Fe3P, which melts as a eutectic 1048 °C (1918 °F) and forms a transient liquid phase. An example of a permanent liquid phase is the infiltration of iron with copper in which 10 to 20 wt% copper remains as a liquid inside the iron skeleton, prior to cooling. The iron-copper infiltrated parts do not shrink. The permanent liquid in the M-2 tool steel causes about 6% linear shrinkage and near-full densification. The process operating in sintering to form necks or bonds between the particles include surface, volume, and grain-boundary diffusion. In the initial stages of sintering, the atoms move over the surfaces of the particles toward the small radius of curvature regions, where the particles contact each other. During sintering, the compact initially shows increases in strength, thermal conductivity, and electrical conductivity. The particle contacts formed in pressing become larger. Later, the strength increases. Finally, toughness, ductility, and density increase with longer sintering time. The structure changes from an interlocked aggregate of powders to what may be thought of as a solid material containing small pores of various sizes and shapes. At longer sintering times, the densification results in the isolation of the pores. They are no longer connected to the surface. To obtain nearly fully density, it is necessary to sinter at temperatures approaching the melting point of the powder. A more detailed sintering theory is found in Volume 7 of the 9th Edition of Metals Handbook, p 309−321. Sintering Techniques. The equipment and techniques used in sintering are diversified; there must be a means of heating the parts in a suitable protective atmosphere and a means of controlling the process variables. Heating and cooling rates, time and temperature of sintering, and composition of the atmosphere are the most critical variables. The dimensions of the finished parts can be affected by the sintering conditions, as well as by the compacting process and the properties of the powder itself. In the sintering of ferrous metals, the atmosphere serves a dual purpose: It both protects the metal from oxidation and helps to control the carbon content of the parts, especially the carbon content at the surface. Commercial ferrous powders, including prealloyed


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01 Sep 2005

steel powders, generally contain very little carbon. The carbon content of the finished part is regulated by the amount of graphite mixed into the powder and the carbon potential of the sintering atmosphere. Particularly in alloy steels, the rate of cooling from the sintering temperature can determine which microconstituents are formed as a result of austenite decomposition, but the rate of cooling generally is slow enough to produce ferritic or pearlitic microstructures. Sintering Equipment. Most iron parts are sintered in mesh belt conveyor furnaces at 1105 to 1120 °C (2020 to 2050 °F) in a mixture of nitrogen and 10% dissociated ammonia. The parts pass through a delubrication zone at 760 to 870 °C (1400 to 1600 °F) and then spend about 15 to 25 min at the high-heat temperature. These furnaces produce 90 to 360 kg/h (200 to 800 lb/h) of parts. They have been improved by the addition of preheat zones that heat rapidly using gas flames to provide an oxidizing environment, which prevents blistering and sooting of parts. Other furnaces have forced atmosphere gas recirculating coolers at the discharge end. These can give finer pearlite spacing or can even form some martensite or bainite, for increased strength. For higher temperatures and improved toughness, parts are now being sintered in large-scale walking beam furnaces at 1230 to 1315 °C (2250 to 2400 °F). Yield can be up to 900 kg/h (2000 lb/h), with good dimensional control and decreased spread in mechanical properties, as well as increases in the mean properties. Mesh belts of SiC were introduced in 1989, allowing sintering at 1315 °C (2400 °F) and higher. Stainless steel, silicon iron, and tool steels are usually sintered in vacuum furnaces. The lubricant is generally removed first in a separate furnace. Such delubrication furnaces have been linked to vacuum furnaces in such a way that trays of parts are automatically delubricated, transferred for sintering, and then transferred for gas pressure quenching. Rapid quenching in gas prevents the pickup of nitrogen in 316 stainless steel and promotes better corrosion resistance. Some furnaces are capable of 1315 to 1540 °C (2400 to 2800 °F). Furnaces with sweep gas systems can remove the binder in the vacuum furnace without the use of a separate atmosphere furnace.

Secondary Operations In many applications, a P/M part that is made by compacting and sintering meets every performance requirements. In other instances, however, the functional requirements (mechanical properties, surface finish, and/or dimensional tolerances) for a part exceed the capabilities of compacted and sintered parts, and one or more secondary operations are required. Some of the common secondary operations are: • • • • • • • • •

Sizing: To tighten dimensional tolerances, usually in the radial direction, relative to the direction of compacting pressure Coining: To change axial dimensions and tolerances Machining: To obtain shapes that cannot be compacted, such as by tapping holes or cutting undercut grooves Forming: To change the shape of the part; can be done hot or cold Re-pressing: To reduce porosity and increase strength and ductility; may be accompanied by resintering Infiltration: To increase strength and decrease porosity Heat treating: To increase hardness and strength Joining: By sinter bonding, staking, brazing, infiltrating, or welding Finishing: Includes deburring, polishing, impregnating, and plating

When the application of a P/M part requires high levels of strength, toughness, or hardness, the mechanical properties can be improved or modified by infiltration, heat treatment, or a secondary mechanical forming operation such as cold (room-temperature) re-pressing or powder forging (hot forming). The effect of these secondary processes on P/M mechanical properties is discussed in later sections of this article.

Designation of P/M Materials Powder metallurgy materials are customarily designated by the specifications or standards to which they are made, such as those listed in Tables 2 and 3 . Comparable standards are published by ASTM, SAE, and MPIF (Metal Powder Industries Federation). Table 2 Compositions of ferrous P/M structural materials Designation (a) Description

MPIF

ASTM

MPIF composition limits and ranges, %(b) SAE

C

Ni

Cu

Fe

Mo

P/M iron

F-0000

B 783

853, Cl 1

0.3 max

...

...

97.7−100

...

P/M steel

F-0005

B 783

853, Cl 2

0.3−0.6

...

...

97.4−99.7

...

P/M steel

F-0008

B 783

853, Cl 3

0.6−1.0

...

...

97.0−99.1

...

P/M copper iron

FC-0200

B 783

...

0.3 max

...

1.5−3.9

93.8−98.5

...

P/M copper steel

FC-0205

B 783

...

0.3−0.6

...

1.5−3.9

93.5−98.2

...


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P/M copper steel

FC-0208

B 783

P/M copper steel

FC-0505

B 783

P/M copper steel

FC-0508

B 783

P/M copper steel

FC-0808

B 783

864, Gr 1, Cl 3

01 Sep 2005

0.6−1.0

...

1.5−3.9

93.1−97.9

...

0.3−0.6

...

4.0−6.0

91.4−95.7

...

864, Gr 2, Cl 3

0.6−1.0

...

4.0−6.0

91.0−95.4

...

864, Gr 3, Cl 3

0.6−1.0

...

6.0−11.0

86.0−93.4

...

8864, Gr 4, Cl 3

0.6−0.9

...

18.0−22.0

75.1 min

...

862

0.3 max

...

...

P/M copper steel

...

...

P/M iron-copper

FC-1000

B 783

9.5−10.5

87.2−90.5

...

P/M prealloyed steel

FL-4205

B 783

...

0.4−0.7

0.35−0.45

...

95.9−98.7

0.50−0.85

P/M prealloyed steel

FL-4605

B 783

...

0.4−0.7

1.70−2.00

...

94.5−97.5

0.40−0.80

P/M iron-nickel

FN-0200

B 783

...

0.3 max

1.0−3.0

2.5 max

92.2−99.0

...

P/M nickel steel

FN-0205

B 783

...

0.3−0.6

1.0−3.0

2.5 max

91.9−98.7

...

P/M nickel steel

FN-0208

B 783

...

0.6−0.9

1.0−3.0

2.5 max

91.6−98.4

...

P/M iron-nickel

FN-0400

B 783

...

0.3 max

3.0−5.5

2.0 max

90.2−97.0

...

P/M nickel steel

FN-0405

B 783

...

0.3−0.6

3.0−5.5

2.0 max

89.9−96.7

...

P/M nickel steel

FN-0408

B 783

...

0.6−0.9

3.0−5.5

2.0 max

89.6−96.4

...

P/M iron-nickel

FN-0700

...

...

0.3 max

6.0−8.0

2.0 max

87.7−94.0

...

P/M nickel steel

FN-0705

...

...

0.3−0.6

6.0−8.0

2.0 max

87.4−93.7

...

P/M nickel steel

FN-0708

...

...

0.6−0.9

6.0−8.0

2.0 max

87.1−93.4

...

P/M infiltrated steel

FX-1000

B 783

...

0−0.3

...

8.0−14.9

82.8−92.0

...


FX-1005

B 783

...

0.3−0.6

...

8.0−14.9

80.5−91.7

...


FX-1008

B 783

...

0.6−1.0

...

8.0−14.9

80.1−91.4

...


FX-2000

B 783

0.3 max

...

15.0−25.0

70.7−85.0

...


FX-2005

B 783

0.3−0.6

...

15.0−25.0

70.4−84.7

...

870 ...

P/M infiltrated steel FX-2008 B 783 872 ... ... 0.6−1.0 15.0−25.0 70.0−84.4 (a) Designations listed are nearest comparable designations; ranges and limits may vary slightly between comparable designations. (b) MPIF standards require that the total amount of all other elements be less than 2.0%, except in infiltrated steels, for which the total amount of other elements must be less than 4.0%

Table 3 Compositions of P/M stainless steels Composition, % MPIF designation

Fe

Cr

Ni

Mn

Si

S

C

P

Mo

N

SS-303N1, N2

rem

17.0−19.00

8.0−13.0

0−2.0

0.−1.0

0.15−0.30

0−0.15

0−0.20

...

0.2−0.6

SS-303L

rem

17.0−19.0

8.0−13.0

0−2.0

0−1.0

0.15−0.30

0−0.03

0−0.20

...

...

SS-304N1, N2

rem

18.0−20.0

8.0−12.0

0−2.0

0−1.0

0−0.03

0−0.08

0−0.045

...

0.2−0.6

SS-304L

rem

18.0−20.0

8.0−12.0

0−2.0

0−1.0

0−0.03

0−0.03

0−0.045

...

...

SS-316N1, N2

rem

16.0−18.0

10.0−14.0

0−2.0

0−1.0

0−0.03

0−0.08

0−0.045

2.0−3.0

0.2−0.6

SS-316L

rem

16.0−18.0

10.0−14.0

0−2.0

0−1.0

0−0.03

0−0.03

0−0.045

2.0−3.0

...

SS-410

rem

11.5−13.0

...

0−1.0

0−1.0

0−0.03

0−0.25

0−0.04

...

0.2−0.6

The MPIF designations for ferrous P/M materials, described in detail in Ref 3, include a prefix of one or more letters (the first of which is F to indicate an iron-base material), four numerals, and a suffix. The second letter in the prefix identifies the principal alloying element (if one is specified); the percentage of that element is indicated by the first two digits. The third and fourth digits indicate the amount of carbon in the compacted and sintered part; the code designation 00 indicates less than 0.3%, 05 indicates 0.3 to 0.6%, and 08 indicates 0.6 to 0.9%. The suffix is used to indicate the minimum 0.2% yield strength of as-sintered parts and the minimum ultimate tensile strength of heat-treated materials in units of 1000 psi (6.894 MPa). The letters HT designate heat treated. Commercially produced iron-base powders often contain controlled amounts of alloying elements other than those specified by any of the designations listed in Table 2 . Manganese and molybdenum may be added to improve strength and the response to heat treatment. Sulfur may be added to enhance machinability. Additions of 0.45 to 0.80% P can improve the toughness of the part and reduce magnetic hysteresis losses. These powders are usually identified by the trade name of the producer even though the amounts of alloy additions are small enough that the designations listed in Table 2 could be applied to the powders. Commercially produced iron-base powders usually contain very little carbon because carbon lowers compressibility and the amount of carbon in the finished part is readily controlled by the amount of admixed graphite and the composition of the


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sintering atmosphere. Commercially produced stainless steels also contain controlled amounts of alloying elements, particularly chromium and nickel. Table 3 lists the designations and compositions of some typical stainless steel P/M materials.

Mechanical and Physical Properties of Sintered Ferrous P/M Materials The mechanical and physical properties of compacted and sintered ferrous powder metallurgy materials depend on many factors. Some factors, such as microstructure, chemical composition, and heat treatment, affect the properties of P/M materials just as they affect the properties of wrought or cast steels. However, these factors themselves are affected by conditions unique to the P/M process, such as type of iron powder, whether alloying additions are prealloyed or admixed, type and amount of lubricant added to the powder, variations in the compacting process, and sintering conditions. For example, prealloyed powders generally produce stronger and tougher parts than do powders mixed from elemental metal powders, assuming that the same compacted density is obtained with each powder. The variations in the P/M process, such as compacting pressure and sintering temperature, also affect the properties of P/M materials. Porosity. The most obvious distinction between the microstructures of wrought and powder metallurgy materials is the porosity often found in powder metallurgy materials. This porosity originates as the spaces between powder particles and persists to some extent through sintering and subsequent secondary operations. The microstructure in Fig. 3 shows an example of a pore in a P/M material. Fig. 3 Atomized iron powder with 0.3% graphite added to yield 0.1 to 0.2% combined carbon (6.7 g/cm3). Pressed at 410 to 480 MPa (30 to 35 tsi) and sintered 30 min at 1120 °C (2050 °F) in dissociated ammonia. White regions are ferrite. Arrows E surround a colony of eutectoid (pearlite). Arrow P points to a pore. 2% nital. 545x

As was shown in Fig. 2 , different base irons of the same density or porosity level have widely differing tensile properties. For that reason, a minimum density is no longer a requirement for a P/M material. P/M materials are specified in terms of the guaranteed minimum 0.2% tensile yield strength for as-sintered materials, and the minimum ultimate tensile strength for heat-treated materials. Nonetheless, for a specific base iron powder and degree of sintering, mechanical properties do increase with density. Under the guaranteed minimum property system, the design engineer selects the tensile properties and chemistry he needs, and the parts fabricator chooses the density, base iron, sintering, and heat treatment to provide the properties. Table 4 shows all the guaranteed minimum tensile properties and related typical properties for the P/M materials that have been standardized by the MPIF. Table 4 Minimum and typical mechanical properties of ferrous P/M materials Minimum strength values (in ksi) are specified by the suffix of the material designation code in the first column of the table. Typical values are given in the remaining columns. 0.2% offset Transverse Ultimate yield rupture Impact Fatigue Elongatio strength strength Elastic modulus strength energy(b) strength(d) Material n in designation code(a)

MP a ksi

MPa

ksi

25 mm (1 in.), %

GPa

106 psi

MPa

ksi

J

ft · lbf

4

3

Apparent hardness(c)

MPa

ksi

48

7

Iron and carbon steel F-000-10(e)

125

18

90

13

1.5

96.5

14.0

248

36

F-0000-15(e)

172

25

125

18

2.5

117

17.0

345

50

8

6

60 HRF

70

10

F-0000-20(e)

262

38

172

25

7.0

141

20.5

655

95

47

35

80 HRF

96

14

F-0005-15(e)

165

24

125

18

80 to 90%. Type 316 can be


Page 1313


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01 Sep 2005

used in contact with sulfuric acid ≤10% at temperatures ≤50 °C (120 °F) if the solutions are aerated; the attack is greater in air-free solutions. Type 317 may be used at temperatures as high as 65 °C (150 °F) with ≤5% concentration. The presence of other materials may markedly change the corrosion rate. As little as 500 to 2000 ppm of cupric ions make it possible to use type 304 in hot solutions of moderate concentration. Other additives may have the opposite effect. Sulfurous acid

Type 304 may be subject to pitting, particularly if some sulfuric acid is present. Type 316 is usable at moderate concentrations and temperatures.

Bases Ammonium hydroxide, sodium hydroxide, caustic solutions

Steels in the 300 series generally have good corrosion resistance at virtually all concentrations and temperatures in weak bases, such as ammonium hydroxide. In stronger bases, such as sodium hydroxide, there may be some attack, cracking, or etching in more concentrated solutions and/or at higher temperatures. Commercial-purity caustic solutions may contain chlorides, which will accentuate any attack and may cause pitting of type 316, as well as type 304.

Organics Acetic acid

Acetic acid is seldom pure in chemical plants but generally includes numerous and varied minor constituents. Type 304 is used for a wide variety of equipment including stills, base heaters, holding tanks, heat exchangers, pipelines, valves, and pumps for concentrations ≤99% at temperatures ≤ ~50 °C (120 °F). Type 304 is also satisfactory⎯if small amounts of turbidity or color pickup can be tolerated⎯for room temperature storage of glacial acetic acid. Types 316 and 317 have the broadest range of usefulness, especially if formic acid is also present or if solutions are unaerated. Type 316 is used for fractionating equipment, for 30−99% concentrations where type 304 cannot be used, for storage vessels, pumps, and process equipment handling glacial acetic acid, which would be discolored by type 304. Type 316 is likewise applicable for parts having temperatures >50 °C (120 °F), for dilute vapors, and for high pressures. Type 317 has somewhat greater corrosion resistance than type 316 under severely corrosive conditions. None of the stainless steels has adequate corrosion resistance to glacial acetic acid at the boiling temperature or at superheated vapor temperatures.

Aldehydes

Type 304 is generally satisfactory.

Amines

Type 316 is usually preferred to type 304.

Cellulose acetate

Type 304 is satisfactory for low temperatures, but type 316 or type 317 is needed for high temperatures.

Formic acids

Type 304 is generally acceptable at moderate temperatures, but type 316 is resistant to all concentrations at temperatures up to boiling.

Esters

With regard to corrosion, esters are comparable to organic acids.

Fatty acids

Type 304 is resistant to fats and fatty acids ≤ ~ 150 °C (300 °F), but type 316 is needed at 150−260 °C (300−500 °F), and type 317, at higher temperatures.

Paint vehicles

Type 316 may be needed if exact color and lack of contamination are important.

Phthalic anhydride

Type 316 is usually used for reactors, fractionating columns, traps, baffles, caps, and piping.

Soaps

Type 304 is used for parts such as spray towers, but type 316 may be preferred for spray nozzles and flake-drying belts to minimize off-color product.

Synthetic detergents

Type 316 is used for preheat, piping, pumps, and reactors in catalytic hydrogenation of fatty acids to give salts of sulfonated high-molecular alcohols.

Tall oil (pulp and paper industry)

Type 304 has only limited use in tall-oil distillation service. High rosin acid streams can be handled by type 316L with a minimum molybdenum content of 2.75%. Type 316 can also be used in the more corrosive high fatty acid streams at temperatures ≤245 °C (475 °F), but type 317 will probably be required at higher temperatures.

Tar

Tar distillation equipment is almost all type 316 because coal tar has a high chloride content; type 304 does not have adequate resistance to pitting.

Urea

Type 316L is generally required.

Type 316 is usually selected for all parts in contact with the product because of its inherent corrosion resistance and greater assurance of product purity. (a) The stainless steels mentioned may be considered for use in the indicated environments. Additional information or corrosion expertise may be necessary prior to use in some environments; for example, some impurities may cause localized corrosion (such as chlorides causing pitting or stress-corrosion cracking of some grades). Source: Ref 3 Pharmaceuticals

General corrosion is often much less serious than localized forms such as stress-corrosion cracking, crevice corrosion in tight


Page 1314



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spaces or under deposits, pitting attack, and intergranular attack in sensitized material such as weld heat-affected zones (HAZ). Such localized corrosion can cause unexpected and sometimes catastrophic failure while most of the structure remains unaffected, and therefore must be considered carefully in the design and selection of the proper grade of stainless steel. Corrosive attack can also be increased dramatically by seemingly minor impurities in the medium that may be difficult to anticipate but that can have major effects, even when present in only parts-per-million concentrations; by heat transfer through the steel to or from the corrosive medium; by contact with dissimilar metallic materials; by stray electrical currents; and by many other subtle factors. At elevated temperatures, attack can be accelerated significantly by seemingly minor changes in atmosphere that affect scaling, sulfidation, or carburization. Despite these complications, a suitable steel can be selected for most applications on the basis of experience, perhaps with assistance from the steel producer. Laboratory corrosion data can be misleading in predicting service performance. Even actual service data have limitations, because similar corrosive media may differ substantially because of slight variations in some of the corrosion factors listed above. For difficult applications, the extensive study of comparative data may be necessary, sometimes followed by pilot plant or in-service testing. More detailed information is available in the section "Corrosion Properties" in this article. Mechanical properties at service temperature are obviously important, but satisfactory performance at other temperatures must be considered also. Thus, a product for arctic service must have suitable properties at subzero temperatures even though steady-state operating temperature may be much higher; room-temperature properties after extended service at elevated temperature can be important for applications such as boilers and jet engines, which are intermittently shut down. Fabrication and Cleaning. Frequently a particular stainless steel is chosen for a fabrication characteristic such as formability or weldability. Even a required or preferred cleaning procedure may dictate the selection of a specific type. For instance, a weldment that is to be cleaned in a medium such as nitric-hydrofluoric acid, which attacks sensitized stainless steel, should be produced from stabilized or low-carbon stainless steel even though sensitization may not affect performance under service conditions. Experience in the use of stainless steels indicates that many factors can affect their corrosion resistance. Some of the more prominent factors are: • • • • • • • • • • • • • • • • • • • • • • •

Chemical composition of the corrosive medium, including impurities Physical state of the medium⎯liquid, gaseous, solid, or combinations thereof Temperature Temperature variations Aeration of the medium Oxygen content of the medium Bacteria content of the medium Ionization of the medium Repeated formation and collapse of bubbles in the medium Relative motion of the medium with respect to the steel Chemical composition of the metal Nature and distribution of microstructural constituents Continuity of exposure of the metal to the medium Surface condition of the metal Stresses in the metal during exposure to the medium Contact of the metal with one or more dissimilar metallic materials Stray electric currents Differences in electric potential Marine growths such as barnacles Sludge deposits on the metal Carbon deposits from heated organic compounds Dust on exposed surfaces Effects of welding, brazing, and soldering

Surface Finish. Other characteristics in the stainless steel selection checklist are vital for some specialized applications but of little concern for many applications. Among these characteristics, surface finish is important more often than any other except corrosion resistance. Stainless steels are sometimes selected because they are available in a variety of attractive finishes. Surface finish selection may be made on the basis of appearance, frictional characteristics, or sanitation. The effect of finish on sanitation sometimes is thought to be simpler than it actually is, and tests of several candidate finishes may be advisable. The selection of finish may in turn influence the selection of the alloy because of differences in availability or durability of the various finishes for


Page 1315



01 Sep 2005

different types. For example, a more corrosion-resistant stainless steel will maintain a bright finish in a corrosive environment that would dull a lower-alloy type. Selection among finishes is described in more detail in this article in the section "Surface Finishing of Stainless Steel."

Product Forms Stainless steels are available in the form of plate, sheet, strip, foil, bar, wire, semifinished products, pipes, tubes, and tubing. Plate Plate is a flat-rolled or forged product more than 250 mm (10 in.) in width and at least 4.76 mm (0.1875 in.) in thickness. Stainless steel plate is produced in most of the types shown in Table 2 . Exceptions include highly alloyed ferritic stainless steels, some of the martensitic stainless steels, and a few of the free-machining grades. Plate is usually produced by hot rolling from slabs that have been directly cast or rolled from ingots and that usually have been conditioned to improve plate surface. Some plate may be produced by direct rolling from ingot. This plate is referred to as sheared plate or sheared mill plate when rolled between horizontal rolls and trimmed on all edges, and as universal plate or universal mill plate when rolled between horizontal and vertical rolls and trimmed only on the ends. Universal plate is sometimes rolled between grooved rolls. Stainless steel plate is generally produced in the annealed condition and is either blast cleaned or pickled. Blast cleaning is generally followed by further cleaning in appropriate acids to remove surface contaminants such as particles of steel picked up from the mill rolls. Plate can be produced with mill edge and uncropped ends. Sheet Sheet is a flat-rolled product in coils or cut lengths at least 610 mm (24 in.) wide and less than 4.76 mm (0.1875 in.) thick. Stainless steel sheet is produced in nearly all types shown in Table 2 except the free-machining and certain martensitic grades. Sheet from the conventional grades is almost exclusively produced on continuous mills. Hand mill production is usually confined to alloys that cannot be produced economically on continuous mills, such as certain high-temperature alloys. The steel is cast in ingots, and the ingots are rolled on a slabbing mill or a blooming mill into slabs or sheet bars. The slabs or sheet bars are then conditioned prior to being hot rolled on a finishing mill. Alternatively, the steel may be continuous cast directly into slabs that are ready for hot rolling on a finishing mill. The current trend worldwide is toward greater production from continuous cast slabs. Sheet produced from slabs on continuous rolling mills is coiled directly off the mill. After they are descaled, these hot bands are cold rolled to the required thickness, and coils off the cold mill are either annealed and descaled or bright annealed. Belt grinding to remove surface defects is frequently required at hot bands or at an intermediate stage of processing. Full coils or lengths cut from coils may then be lightly cold rolled on either dull or bright rolls to produce the required finish. Sheet may be shipped in coils, or cut sheets may be produced by shearing lengths from a coil and flattening them by roller leveling or stretcher leveling. Sheet produced on hand mills from sheet bars is rolled in lengths and then annealed and descaled. It may be subjected to additional operations, including cold reduction, annealing, descaling, light cold rolling for finish, or flattening. A specified minimum tensile strength, minimum yield strength, or hardness level higher than that normally obtained on sheet in the annealed condition, or a combination thereof, can be attained by controlled cold rolling. Sheet made of chromium-nickel stainless steel (often type 301) or of chromium-nickel-manganese stainless steel (often type 201) is produced in the following cold-rolled tempers: Minimum tensile strength Temper

Minimum yield strength

MPa

ksi

MPa

1

860

125

515

75

1

1035

150

760

110

3

=4 hard

1205

175

930

135

Full hard

1275

185

965

140

=4 hard =2 hard

ksi

Strip Strip is a flat-rolled product, in coils or cut lengths, less than 610 mm (24 in.) wide and 0.13 to 4.76 mm (0.005 to 0.1875 in.) thick. Cold finished material 0.13 mm (0.005 in.) thick and less than 610 mm (24 in.) wide fits the definitions of both strip and foil and may be referred to by either term. Cold-rolled stainless steel strip is manufactured from hot-rolled, annealed, and pickled strip (or from slit sheet) by rolling between polished rolls. Depending on the desired thickness, various numbers of cold-rolling passes through the mill are required for effecting the necessary reduction and securing the desired surface characteristics and mechanical properties.


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Hot-rolled stainless steel strip is a semifinished product obtained by hot-rolling slabs or billets and is produced for conversion to finished strip by cold rolling. Heat Treatment. Strip of all types of stainless steel is usually either annealed or annealed and skin passed, depending on requirements. When severe forming, bending, and drawing operations are involved, it is recommended that such requirements be indicated so that the producer will have all the information necessary to ensure that he supplies the proper type and condition. When stretcher strains are objectionable in ferritic stainless steels such as type 430, they can be minimized by specifying a No. 2 finish. Cold-rolled strip in types 410, 414, 416, 420, 431, 440A, 440B, and 440C can be produced in the hardened and tempered condition. Strip made of chromium-nickel stainless steel (often type 301) or of chromium-nickel-manganese stainless steel (often type 201) is produced in the same cold-rolled tempers in which sheet is produced. For strip, edge condition is often important⎯more important than it usually is for sheet. Strip can be furnished with various edge specifications: • • • •

Mill edge (as produced, condition unspecified) No. 1 edge (edge rolled, rounded, or square) No. 3 edge (as slit) No. 5 edge (square edge produced by rolling or filing after slitting)

Mill edge is the least expensive edge condition, and is adequate for many purposes. No. 1 edge provides improved width tolerance over mill edge plus a cold-rolled edge condition; rounded edges are preferred for applications requiring the lowest degree of stress concentration at corners. No. 3 and No. 5 edges give progressively better width tolerance and squareness over No. 1 edge. Foil Foil is a flat-rolled product, in coil form, up to 0.13 mm (0.005 in.) thick and less than 610 mm (24 in.) wide. Foil is produced in slit widths with edge conditions corresponding to No. 3 and No. 5 edge conditions for strip. Foil is made from types 201, 202, 301, 302, 304, 304L, 305, 316, 316L, 321, 347, 430, and 442, as well as from certain proprietary alloys. The finishes, tolerances, and mechanical properties of foil differ from those of strip because of limitations associated with the way in which foil is manufactured. Nomenclature for finishes, and for width and thickness tolerances, vary among producers. Finishes for foil are described by the finishing operations employed in their manufacture. However, each finish in itself is a category of finishes, with variations in appearance and smoothness that depend on composition, thickness, and method of manufacture. Chromium-nickel and chromium-nickel-manganese stainless steels have a characteristic appearance different from that of straight chromium types for corresponding finish designations. Mechanical Properties. In general, mechanical properties of foil vary with thickness. Tensile strength is increased somewhat, and ductility is lowered, by a decrease in thickness. Bar Bar is a product supplied in straight lengths; it is either hot or cold finished and is available in various shapes, sizes, and surface finishes. This category includes small shapes whose dimensions do not exceed 75 mm (3 in.) and, secnd, hot-rolled flat stock at least 3.2 mm (0.125 in.) thick and up to 250 mm (10 in.) wide. Hot-finished bar is commonly produced by hot rolling, forging, or pressing ingots to blooms or billets of intermediate size, which are subsequently hot rolled, forged, or extruded to final dimensions. Whether rolling, forging, or extrusion is selected as the finishing method depends on several factors, including composition and final size. Following hot rolling or forging, hot-finished bar may be subjected to various operations, including: • • • •

Annealing or other heat treatment Descaling by pickling, blast cleaning, or other methods Surface conditioning by grinding or rough turning Machine straightening

Cold-finished bar is produced from hot-finished bar or rod by additional operations such as cold rolling or cold drawing, which result in the close control of dimensions, a smooth surface finish, and higher tensile and yield strengths. Sizes and shapes of cold reduced stock classified as bar are essentially the same as for hot-finished bar, except that all cold reduced flat stock less than 4.76 mm (0.1875 in.) thick and over 9.5 mm (0.375 in.) wide is classified as strip. Cold-finished round bar is commonly machine straightened; afterward, it can be centerless ground or centerless ground and polished. Centerless grinding and polishing do not alter the mechanical properties of cold-finished bar and are used only to improve surface finish or provide closer tolerances. Some increase in hardness, more marked at the surface and particularly in 2xx and 3xx stainless steels, results from machine straightening. The amount of increase varies chiefly with composition, size, and


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amount of cold work necessary to straighten the bar. Cold-finished bars that are square, flat, hexagonal, octagonal, or of certain special shapes are produced from hot-finished bars by cold drawing or cold rolling. When cold-finished bar is required to have high strength and hardness, it is cold drawn or heat treated, depending on composition, section size, and required properties. Round sections can be subsequently centerless ground or centerless ground and polished. Free-machining wire is a bar commodity used for making parts in automatic screw machines or other types of machining equipment. The principal types used are 303, 303Se, 416, 416Se, 420F, 430F, and 430FSe. Free-machining wire is commonly produced with a cold drawn or centerless ground finish and with selected hardnesses, depending on the machining operation involved. Structural Shapes. Hot-rolled, bar-size structural shapes are produced in angles, channels, tees, and zees. They can be purchased in various conditions: • • • • •

Hot rolled Hot rolled and annealed Hot rolled, annealed, and blast cleaned Hot rolled, annealed, and chemically cleaned Hot rolled, annealed, blast cleaned, and chemically cleaned

Wire Wire is a coiled product derived by cold finishing hot-rolled and annealed rod. Cold finishing imparts excellent dimensional accuracy, good surface smoothness, a fine finish, and specific mechanical properties. Wire is produced in several tempers and finishes. Wire is customarily referred to as round wire when the contour is completely cylindrical and as shape wire when the contour is other than cylindrical. For example, wires that are half round, half oval, oval, square, rectangular, hexagonal, octagonal, or triangular in cross section are all referred to as shape wire. Shape wire is cold finished either by drawing or by a combination of drawing and rolling. In the production of wire, rod (which is a coiled hot-rolled product approximately round in cross section) is drawn through the tapered hole of a die or a series of dies. The smallest size of hot-rolled rod commonly made is 5.5 mm (0.218 in.). Rod smaller than this is produced by cold work, the number of dies employed depending on the finished diameter required. Round stainless steel wire is commonly produced within the approximate size range 0.08 to 15.9 mm (0.003 to 0.625 in.). Shape wire, except cold-finished flat wire, is commonly produced within the approximate size range of 1.12 to 12.7 mm (0.044 to 0.500 in.), although the particular shape governs the specific sizes that can be produced. Tempers of Wire. There are four classifications of wire temper: annealed-temper, soft-temper, intermediate-temper, and spring-temper. Annealed temper describes soft wire that has undergone no further cold drawing after the last annealing treatment. Wire in this temper is made by annealing in open-fired furnaces or molten salt, and annealing ordinarily is followed by pickling that produces a clean, gray, matte finish. It is also made with a bright finish by annealing in a protective atmosphere and sometimes is described as bright annealed wire. Soft-temper wire is given a single light draft following the final annealing operation and generally is produced to a defined upper limit of tensile strength or hardness. Wire in this temper is produced with various dry-drawn finishes, including lime soap, lead, copper, and oxide. It may also be given a bright finish produced by oil or grease drawing. Intermediate-temper wire is drawn one or more drafts after annealing as required to produce a specific minimum strength or hardness. The properties of this wire can vary between the properties of soft-temper wire and properties approaching those of spring-temper wire. Intermediate-temper wire is usually produced with one of the dry-drawn finishes. Spring-temper wire is drawn several drafts as required to produce high tensile strengths. Special Wire Commodities. There are many classes of stainless steel wire that have been developed for specific components or for particular applications. The unique properties of each of these individual wire commodities are developed by employing a particular combination of composition, steel quality, process heat treatment, and cold drawing practice. The details of manufacture may vary slightly from one wire manufacturer to another, but the finished wire will fulfill the specified requirements. Cold-heading wire is produced in any of the various types of stainless steel. In all instances, cold-heading wire is subjected to special testing and inspection to ensure satisfactory performance in cold-heading and cold-forging operations. Of the chromium-nickel group, types 305 and 302Cu are used for cold-heading wire and generally are necessary for severe upsetting. Other grades commonly cold formed include 304, 316, 321, 347, and 384. Of the 4xx series, types 410, 420, 430, and 431 are used for a variety of cold-headed products. Types 430 and 410 are commonly used for severe upsetting and for recessed-head screws and bolts. Types 416, 416Se, 430F, and 430FSe are intended primarily for free cutting and are not recommended for cold heading. Cold-heading wire is manufactured using a closely controlled annealing treatment that produces optimum softness and still


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permits a very light finishing draft after pickling. The purposes of the finishing draft are to provide a lubricating coating that will aid the cold-heading operation and to produce a kink-free wire coil having more uniform dimensions. Cold-heading wire is produced with a variety of finishes, all of which have the function of providing proper lubrication in the header dies. The finish or coating should be suitably adherent to prevent galling and excessively rapid die wear. A copper coating, which is applied after the annealing treatment and just prior to the finishing draft, is available; the copper-coated wire is then lime coated and drawn, using soap as the drawing lubricant. Coatings of lime and soap or of oxide and soap are also employed. Spring wire is drawn from annealed rod and is subjected to mill tests and inspection that ensures the quality required for extension and compression springs. The types of stainless steel of which spring wire is commonly produced include 302, 304, and, for additional corrosion resistance, 316, and UNS N08020. Spring wire in large sizes can be furnished in a variety of finishes, such as dry-drawn lead, copper, lime and soap, and oxide and soap. Fine sizes are usually wet drawn, although they can be dry drawn. Tensile strength ranges or minimums for types 302, 304, 305, and 316 spring wire in various sizes are given in Table 6 . Table 6 Room-temperature tensile strength of stainless steel spring wire Diameter

Tensile strength

mm

in.

MPa

ksi

Types 302 and 304 ≤0.23

≤0.009

2241−2448

325−355

>0.23−0.25

>0.009−0.010

2206−2413

320−350

>0.25−0.28

>0.010−0.011

2192−2399

318−348

>0.28−0.30

>0.011−0.012

2179−2385

316−346

>0.30−0.33

>0.012−0.013

2165−2372

314−344

>0.33−0.36

>0.013−0.014

2151−2358

312−342

>0.36−0.38

>0.014−0.015

2137−2344

310−340

>0.38−0.41

>0.015−0.016

2124−2330

308−338

>0.41−0.43

>0.016−0.017

2110−2317

306−336

>0.43−0.46

>0.017−0.018

2096−2303

304−334

>0.46−0.51

>0.018−0.020

2068−2275

300−330

>0.51−0.56

>0.020−0.022

2041−2248

296−326

>0.56−0.61

>0.022−0.024

2013−2220

292−322

>0.61−0.66

>0.024−0.026

2006−2206

291−320

>0.66−0.71

>0.026−0.028

1993−2192

289−318

>0.71−0.79

>0.028−0.031

1965−2172

285−315

>0.79−0.86

>0.031−0.034

1944−2137

282−310

>0.86−0.94

>0.034−0.037

1930−2124

280−308

>0.94−1.04

>0.037−0.041

1896−2096

275−304

>1.04−1.14

>0.041−0.045

1875−2068

272−300

>1.14−1.27

>0.045−0.050

1841−2034

267−295

>1.27−1.37

>0.050−0.054

1827−2020

265−293

>1.37−1.47

>0.054−0.058

1800−1993

261−289

>1.47−1.60

>0.058−0.063

1779−1965

258−285

>1.60−1.78

>0.063−0.070

1737−1937

252−281

>1.78−1.90

>0.070−0.075

1724−1917

250−278

>1.90−2.03

>0.075−0.080

1696−1896

246−275

>2.03−2.21

>0.080−0.087

1668−1868

242−271

>2.21−2.41

>0.087−0.095

1641−1848

238−268

>2.41−2.67

>0.095−0.105

1600−1806

232−262

>2.67−2.92

>0.105−0.115

1565−1772

227−257

>2.92−3.18

>0.115−0.125

1531−1744

222−253


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01 Sep 2005

>3.18−3.43

>0.125−0.135

1496−1710

217−248

>3.43−3.76

>0.135−0.148

1448−1662

210−241

>3.76−4.12

>0.148−0.162

1413−1620

205−235

>4.12−4.50

>0.162−0.177

1365−1572

198−228

>4.50−4.88

>0.177−0.192

1338−1551

194−225

>4.88−5.26

>0.192−0.207

1296−1517

188−220

>5.26−5.72

>0.207−0.225

1255−1475

182−214

>5.72−6.35

>0.225−0.250

1207−1413

175−205

>6.35−7.06

>0.250−0.278

1158−1365

168−198

>7.06−7.77

>0.278−0.306

1110−1324

161−192

>7.77−8.41

>0.306−0.331

1069−1282

155−186

>8.41−9.20

>0.331−0.362

1020−1241

148−180

>9.20−10.01

>0.362−0.394

979−1193

142−173

>10.01−11.12

>0.394−0.438

931−1138

135−165

>11.12−12.70

>0.438−0.500

862−1069

125−155

≤0.010

1689−1896

245−275

>0.25−0.38

>0.010−0.015

1655−1862

240−270

>0.38−1.04

>0.015−0.041

1620−1827

235−265

>1.04−1.19

>0.041−0.047

1586−1723

230−260

>1.19−1.37

>0.047−0.054

1551−1758

225−255

>1.37−1.58

>0.054−0.062

1517−1724

220−250

>1.58−1.85

>0.062−0.072

1482−1689

215−245

>1.85−2.03

>0.072−0.080

1448−1655

210−240

>2.03−2.34

>0.080−0.092

1413−1620

205−235

>2.34−2.67

>0.092−0.105

1379−1586

200−230

>2.67−3.05

>0.105−0.120

1344−1551

195−225

>3.05−3.76

>0.120−0.148

1276−1482

185−215

>3.76−4.22

>0.148−0.166

1241−1448

180−210

>4.22−4.50

>0.166−0.177

1172−1379

170−200

>4.50−5.26

>0.177−0.207

1103−1310

160−190

>5.26−5.72

>0.207−0.225

1069−1276

155−185

>5.72−6.35

>0.225−0.250

1034−1241

150−180

>6.35−7.92

>0.250−0.312

931−1138

135−165

>7.92−12.68

>0.312−0.499

793−1000

115−145

Types 305 and 316 ≤0.25

>12.68

>0.499

Consult producer

The torsional modulus for stainless steel spring wire may range from 59 to 76 GPa (8.5 to 11 × 106 psi), depending on alloy and wire size. Magnetic permeability is extremely low compared to that of carbon steel wire. Springs made from stainless steel wire retain their physical and mechanical properties at temperatures up to about 315 °C (600 °F). Rope wire is used to make rope, cable, and cord for a variety of uses, such as aircraft control cable, marine rope, elevator cable, slings, and anchor cable. Because of special requirements for fatigue strength, rope wire is produced from specially selected and processed material. Rope wire is made of type 302 or type 304 unless a higher level of corrosion resistance is required, in which case type 316 is generally selected. Special nonmagnetic characteristics may be required, which necessitate the selection of grades that have little or no ferrite or martensite in the microstructure and the use of special drawing practices to limit or avoid deformation-induced transformation to martensite. Tensile properties of regular rope wire are slightly lower than those of stainless steel spring wire. Finishes for rope wire vary from a gray matte finish to a bright finish and include a series of bright to dark soap finishes. Soap finishes afford some lubrication that facilitates laying up of rope and also to some extent aids in-service use.


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Weaving wire is used in the weaving of screens for many different applications in coal mines, sand-and-gravel pits, paper mills, chemical plants, dairy plants, oil refineries, and food-processing plants. Annealing and final drawing must be carefully controlled to maintain uniform temper and finish throughout each coil or spool. Because weaving wire must be ductile, it is usually furnished in the annealed temper with a bright annealed finish, or in the soft temper with either a lime-soap finish or an oil- or grease-drawn finish. Most types of stainless steel are available in weaving wire; the most widely used types are 302, 304, 309, 310, 316, 410, and 430. Annealed wire in the 3xx series commonly has a tensile strength of 655 to 860 MPa (95 to 125 ksi) and an elongation (in 50 mm, or 2 in.) of 35 to 60%. Soft-temper wire, which is commonly specified for sizes over 0.75 mm (0.030 in.), averages 860 to 1035 MPa (125 to 150 ksi) in tensile strength and exhibits 15 to 40% elongation. For annealed wire in types 410 and 430, tensile strength averages 495 to 585 MPa (72 to 85 ksi), and elongation averages 17 to 23%. Armature binding wire is produced in types 302 or 304 stainless steel of a composition that is balanced to produce high tensile and yield strengths and low magnetic permeability. Minimum tensile strength of 1515 MPa (220 ksi), minimum yield strength (0.2% offset) of 1170 MPa (170 ksi), and maximum permeability of 4.0 at 16 kA · m−1 (200 oersteds) are usually specified. The wire must be strong enough to withstand the centrifugal forces encountered in use, yet ductile enough to withstand being bent sharply back on itself without cracking when a hook is formed to hold the armature wire during the binding operation. Armature binding wire is furnished on spools and has a smooth, tightly adherent tinned coating that facilitates soldering. Slide forming wire is produced in all standard types, particularly in types 302, 304, 316, 410, and 430. It can be produced in any temper suitable for forming any of the numerous shapes made on slide-type wire forming machines. Wool wire is designed for the production of wool by shredding. It is commonly furnished in an intermediate temper and produced to rigid standards so that it will perform satisfactorily in the wool-cutting operation. Wool wire usually is made from type 430 and has a lime-soap finish. Reed wire is high-quality wire produced for the manufacture of dents for reeds that, once assembled, are used in weaving textiles and other products. Dents are made by rolling the round reed wire into a flat section, and then machining and polishing the edges to a very smooth and accurate contour before cutting the wire into individual dents. Accuracy in size and shape are necessary because of the various processes that the wire must undergo. Reed wire is usually made from type 430 in an intermediate temper that must be uniform in properties throughout each coil and each shipment. The finish also must be uniform and bright. Lashing wire is designed for lashing electric power transmission lines to support cables. Lashing wire is usually made from type 430. It is furnished in the annealed temper with a bright finish and has a maximum tensile strength of 655 MPa (95 ksi) and minimum elongation of 17% in 255 mm (10 in.). It is normally furnished on coreless spools. Cotter pin wire is approximately half-round wire designed for fabricating cotter pins. It is generally produced by rolling round wire between power-driven rolls, by drawing it between power-driven rolls, or by drawing it through a die or Turk's-head roll. To facilitate the spreading of the cotter pin ends, it is desirable that the flat side of the wire have a small radius rather than sharp corners at the edges. Cotter pin wire is commonly furnished in vibrated or hank-wound coils with the flat side of the wire facing inward. Ordinarily it is produced in the soft temper to prevent undesirable springback in the legs of formed cotter pins. Usually it is furnished with a bright finish, but it is also available with a metallic coating. Stainless welding wire is available for many grades to provide good weldability with optimized mechanical properties and corrosion resistance of the weldment. For example, the weldability of austenitic stainless steels is enhanced by controlling unwanted residual elements or balancing the wire composition to provide a small amount of ferrite in the as-deposited weld metal. Also, the composition of duplex stainless weld wire is generally controlled to produce levels of austenite and ferrite in the weld metal that will optimize mechanical properties and corrosion resistance. Stainless steel weld wire is produced in layer-level wound spools, straight lengths (both included in the American Welding Society AWS A5.9) and coated electrodes (AWS A5.4). Semifinished Products Blooms, billets, and slabs are hot rolled, hot forged, or hot pressed to approximate cross-sectional dimensions and generally have rounded corners. Round billets are also produced, typically for extrusion or closed-die forging. These semifinished products, as well as tube rounds, are produced in random lengths or are cut to specified lengths or to specified weights. There are no invariable criteria for distinguishing between the terms bloom and billet, and often they are used interchangeably. Dimensions. The nominal cross-sectional dimensions of blooms, billets, and slabs are designated in inches and fractions of an inch. The size ranges commonly listed as hot-rolled stainless steel blooms, billets, and slabs include square sections 100 × 100 mm (4 × 4 in.) and larger, and rectangular sections at least 10,300 mm2 (16 in.2) in cross-sectional area. Stainless Types. Blooms, billets, and slabs made of 4xx stainless steels that are highly hardenable (types 414, 420, 420F, 422, 431, 440A, 440B, and 440C) are annealed before shipment to prevent cracking. Other hardenable types, such as 403, 410, 416, and 416Se, also may be furnished in the annealed condition, depending on composition and size. Processing. In general practice, blooms, billets, and slabs are cut to length by hot shearing. Hot sawing and flame cutting are also used. When the end distortion or burrs normally encountered in regular mill cutting are not acceptable, ends can be prepared for subsequent operations by any method that does not leave distortion or burrs. Usually, this is grinding. Blooms, billets, tube rounds, and slabs are surface conditioned by grinding or turning prior to being processed by hot rolling, hot forging, hot extruding, or hot piercing. Material can be tested by ultrasonic and macroetching techniques in the as-worked condition;


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however, a more critical evaluation is possible after the material has been conditioned. At the time an order is placed, producer and customer should come to an agreement regarding the manner in which testing or inspection is to be conducted and results interpreted. Pipe, Tubes, and Tubing Pipe, tubes, and tubing are hollow products made either by piercing rounds or by rolling and welding strip. They are used for conveying gases, liquids, and solids, and for various mechanical and structural purposes. (Cylindrical forms intended for use as containers for storage and shipping purposes and products cast to tubular shape are not included in this category.) The number of terms used in describing sizes and other characteristics of stainless steel tubular products has grown with the industry, and in some cases terms may be difficult to define or to distinguish from one another. For example, the terms pipe, tubes, and tubing are distinguished from one another only by general use, not by clear-cut rules. Pipe is distinguished from tubes chiefly by the fact that it is commonly produced in relatively few standard sizes. Tubing is generally made to more exacting specifications than either pipe or tubes, regarding dimensions, finish, chemical composition, and mechanical properties. Stainless steel tubular products are classified according to intended service, as described in the following paragraphs and tabular matter. Stainless Steel Tubing for General Corrosion-Resisting Service. Straight chromium (ferritic or martensitic) types are produced in the annealed or heat-treated condition, and chromium-nickel (austenitic) types are produced in the annealed or cold-worked condition. Austenitic types are inherently tougher and more ductile than ferritic types for similar material conditions or tempers. ASTM specifications A 268 and A 269 apply to stainless steel tubing for general service: A 268 applies to ferritic grades, and A 269, to austenitic grades. Most ferritic grades are also covered by ASME SA268, which sets forth the same material requirements as does ASTM A 268. Stainless steel pressure pipe is made from straight chromium and chromium-nickel types and is governed by the specifications: Specifications ASTM

ASME

A 312

SA312

Seamless and welded pipe

A 358

SA358

Electric fusion welded pipe for high-temperature service

A 376

SA376

Seamless pipe for high-temperature central-station service

A 409

...

A 790

SA790

Description

Large-diameter welded pipe for corrosion or high-temperature service Seamless and welded ferritic/austenitic stainless steel pipe

Stainless steel pressure tubes include boiler, superheater, condenser, and heat-exchanger tubes, which commonly are manufactured from chromium-nickel types; requirements are set forth in the specifications: Specifications Description

ASTM

ASME

A 213

SA213

Ferritic and austenitic alloy seamless tubes for boilers, superheaters, and heat exchangers

A 249

SA249

Austenitic alloy welded tubes for boilers, superheaters, heat exchangers, and condensers

A 271

SA271

Austenitic alloy seamless still tubes for refinery service

A 498

...

A 688

SA688

Welded austenitic stainless steel feedwater heater tubes

A 789

SA789

Seamless and welded ferritic/austenitic stainless steel tubing

Ferritic and austenitic alloy seamless and welded tubes with integral fins

Stainless steel sanitary tubing is used extensively in the dairy and food industries, where cleanliness and exceptional corrosion resistance are important surface characteristics. In many instances, even the slight amounts of corrosion that result in tarnishing or in release of a few ppm of metallic ions into the process stream are objectionable. Sanitary tubing may be polished on the outside or the inside, or both, to provide smooth, easily cleanable surfaces. Special finishes and close dimensional tolerances for special fittings are sometimes required. ASTM A 270 is in common use for this tubing. Stainless steel mechanical tubing is produced in round, square, rectangular, and special-shape cross sections. It is used for many different applications, most of which do not require the tubing to be pressurized. Mechanical tubing is used for bushings; small cylinders; bearing parts; fittings; various types of hollow, cylindrical or ringlike formed parts; and structural members such as furniture frames, machinery frames, and architectural members. ASTM A 511 and A 554 apply to seamless and welded mechanical tubing, respectively. Stainless steel aircraft tubing, produced from various chromium-nickel types, has many structural and hydraulic applications in aircraft construction because of its high resistance to both heat and corrosion. Work-hardened tubing can be used


Page 1322



01 Sep 2005

in high-strength applications, but it is not recommended for parts that may be exposed to certain corrosive substances or to certain combinations of corrosive static or fluctuating stress. Low-carbon types or compositions stabilized by titanium or by niobium with or without tantalum are commonly used when welding is to be done without subsequent heat treatment. Aircraft tubing is made to close tolerances and with special surface finishes, special mechanical properties, and stringent requirements for testing and inspection. It is used for structural components of aircraft fuselages, engine mounts, engine oil lines, landing gear components, and engine parts and is finding increasing application in parts for hydraulic, fuel-injection, exhaust, and heating systems. Aircraft structural tubing is both seamless and welded stainless steel tubing in sizes larger than those referred to as aircraft tubing. It is commonly used in exhaust systems (including stacks), cross headers, collector rings, engine parts, heaters, and pressurizers. Sometimes, stainless steel aircraft structural tubing is produced especially for parts that are to be machined. Stabilized types are used for welded and brazed structures. Seamless and welded stainless steel aircraft structural tubing is made in sizes ranging from 1.6 to 125 mm (1=16to 15 in.) in outside diameter and from 0.25 to 6.35 mm (0.010 to 0.250 in.) in wall thickness. It is ordinarily produced to the federal and Aerospace Material Specification (AMS) specifications listed below. However, because the U.S. government has embarked on a program of replacing military (MIL) specifications with AMS and ASTM specifications, the MIL specifications listed may no longer apply. UNS number, composition, and condition

Specification Seamless tubing AMS 5560

S30400; 19Cr-9Ni; annealed

AMS 5561

S21900; 21Cr-6Ni-9Mn; annealed

AMS 5570

S32100; 18Cr-11Ni (Ti stabilized); annealed

AMS 5571

S34700; 18Cr-11Ni (Nb + Ta stabilized); annealed

AMS 5572


AMS 5573

S31600; 17Cr-12.5Ni-2.5Mo; annealed

AMS 5574

S30908; 23Cr-13.5Ni; annealed

AMS 5578

S45500; 12.5Cr-8.5Ni-0.03 (Nb + Ta)-1.1Ti-2.0Cu; annealed

Welded tubing MIL-T-6737

18-8 (stabilized); annealed

AMS 5565


AMS 5575

S34700; 18Cr-11Ni (Nb + Ta stabilized); annealed

AMS 5576

S32100; 18Cr-10Ni (Ti stabilized); annealed

AMS 5577


Seamless and welded tubing MIL-T-5695

18-8; hardened (cold worked)

MIL-T-8506

S30400; 18-8; annealed

MIL-T-8686

18-8 (stabilized); annealed

Aircraft Hydraulic-Line Tubing. Stainless steel tubing is used widely in aircraft and aerospace vehicles for fuel-injection lines and hydraulic systems. Most of the tubing used for such applications is relatively small; types 304, 304L, 321, 347, and 21-6-9 are most often specified. Aircraft hydraulic-line tubing must have high strength, high ductility, high fatigue resistance, high corrosion resistance, and good cold-working qualities. The ability to be flared for use with standard flare fittings, the ability to be bent without excessive distortion or fracture, and cleanliness of the inside surface are important requirements. Stainless steel aircraft hydraulic-line tubing is produced in either the annealed or the cold-worked ( 1=8hard) condition. The 1 =8hard temper is used wherever possible to save weight. Specifications for stainless steel aircraft hydraulic-line tubing, either seamless or welded, are: Specification

UNS number, or type, and condition

MIL-T-6845

S30400; 1=8 hard

MIL-T-8504

S30400; annealed

MIL-T-8808

321 or 347; annealed

AMS 5556

S34700; annealed

AMS 5557

S32100; annealed

AMS 5560

S30400; annealed


Page 1323



01 Sep 2005

S30400; 1=8 hard

AMS 5566

Tensile Properties Mechanical properties of most stainless steels, especially ductility and toughness, are higher than the same properties of carbon steels. Strength and hardness can be raised by cold work for ferritic and austenitic types, and by heat treatment for precipitation-hardening and martensitic types. Certain ferritic stainless steels can also be hardened slightly by heat treatment. Austenitic Types. Basic room-temperature properties of standard austenitic stainless steels and of several nonstandard austenitic stainless steels are given in Tables 7 and 8 . Additional specifications for austenitic stainless steels include: Product form

ASTM specification

Wire

A 313, A 492, A 493, A 555, B 471, B 475

Bar and wire

B 649

Wire, bar, shapes

A 479

Billet and bar

A 314, B 472

Flanges, fittings, and/or valves, and so on

A 182, A 403, B 462, A 403

Bolting

A 193

Nuts

A 194

Table 7 Minimum room-temperature mechanical properties of austenitic stainless steels Tensile strength Product form(a)

Condition

0.2% yield strength

Elonga Reductio n tion, in area, Hardness, % % HRB

MPa

ksi

MPa

ksi

620

90

205

30

40

...

95 max

ASTM specification

Type 301 (UNS S30100) B

Annealed

B, P, Sh, St

Annealed

515

75

205

30

40

...

92 max

A 167

B, P, Sh, St

1

860

125

515

75

25

...

...

A 666

B, P, Sh, St

1

1030

150

760

110

18

...

...

A 666

B, P, Sh, St

3

=4 hard

1210

175

930

135

12

...

...

A 666

B, P, Sh, St

Full hard

1280

185

965

140

9

...

...

A 666

=4 hard =2 hard

A 666

Type 302 (UNS S30200) B, F

Hot finished and annealed

515

75

205

30

40

50

...

A 276, A 473

B

Cold finished(b) and annealed

620

90

310

45

30

40

...

A 276

B

Cold finished(c) and annealed

515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

P, Sh, St

Annealed

515

75

205

30

40

...

92 max

B, P, Sh, St

High tensile, 1=4 hard

860

125

515

75

10

...

...

A 666

B, P, Sh, St

High tensile, 1=2 hard

1030

150

760

110

10

...

...

A 666

3

A 167, A 240, A 666

B, P, Sh, St

High tensile, =4 hard

1205

175

930

135

6

...

...

A 666

B, P, Sh, St

Full hard

1275

185

965

140

4

...

...

A 666

Type 302B (UNS S30215) B, F


515

75

205

30

40

50

...

A 276, A 473

B


620

90

310

45

30

40

...

A 276

B


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580


Page 1324


P, Sh, St

Annealed


01 Sep 2005

515

75

205

30

40

...

95 max

A 167

Type 302Cu (UNS S30430) W(e)

Annealed

550

80

...

...

...

...

...

A 493

W(e)

Lightly drafted

585

85

...

...

...

...

...

A 493

205

30

40

50

...

A 473

Types 303 (UNS S30300) and 303Se (UNS S30323) F

Annealed

515

75

W

Annealed

585−860

85−125

...

...

...

...

...

A 581

W

Cold worked

790−1000 115−145

...

...

...

...

...

A 581

Type 304 (UNS S30400) B, F(f)


515

75

205

30

40

50

...

A 276, A 473

B


620

90

310

45

30

40

...

A 276

B


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

P, Sh, St

Annealed

515

75

205

30

40

...

92 max

A 167

B, P, Sh, St

1

690

100

380

55

35

...

...

A 666

B, P, Sh, St

1

860

125

515

75

10

...

...

A 666

B, P, Sh, St

1

1035

150

760

110

7

...

...

A 666

450

65

170

25

40

50

...

A 473

=8 hard =4 hard =2 hard

Type 304L (UNS S30403) F

Annealed

B


480

70

170

25

40

50

...

A 276

B


620

90

310

45

30

40

...

A 276

B


480

70

170

25

30

40

...

A 276

W

Annealed

480

70

170

25

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

P, Sh, St

Annealed

480

70

170

25

40

...

88 max

A 167, A 240

P, Sh, St grade A Annealed

515

75

205

30

27

...

95 max

A 887

P, Sh, St grade B Annealed

515

75

205

30

16

...

95 max

A 887

Type 304B4 (UNS S30424)

Type 305 (UNS S30500) B, F


515

75

205

30

40

50

...

A 276, A 473

B


260

90

310

45

30

40

...

A 276

B


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

P, Sh, St

Annealed

480

70

170

25

40

...

88 max

A 167

B, W

High tensile(d)

1690

245

...

...

...

...

...

540

78

240

35

40

...

...

...

Cronifer 18-15 LCSi (UNS S30600) P, Sh, St

Annealed

A 167, A 240

Type 308 (UNS S30800) B, F


515

75

205

30

40

50

...

A 276, A 473

B


620

90

310

45

30

40

...

A 276

B


515

75

205

30

30

40

...

A 276


Page 1325



01 Sep 2005

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

P, Sh, St

Annealed

515

75

205

30

40

...

88 max

A 167

Types 309 (UNS S30900), 309S (UNS S30908), 310 (UNS S31000) and 310S (UNS S31008) B, F


515

75

205

30

40

50

...

A 276, A 473

B


620

90

310

45

30

40

...

A 276

B


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

P, Sh, St

Annealed

515

75

205

30

40

...

95 max

A 167

310Cb (UNS S31040) P, Sh, St

Annealed

515

75

205

30

40

...

95

A 167, A 240

B, Shapes


515

75

205

30

40

50

...

A 276

B, Shapes


620

90

310

45

30

40

...

A 276

B, Shapes


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

Type 314 (UNS S31400) B, F


515

75

205

30

40

50

...

A 276, A 473

B


620

90

310

45

30

40

...

A 276

B


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

Type 316 (UNS S31600) B, F(f)


515

75

205

30

40

50

...

A 276, A 473

B


620

90

310

45

30

40

...

A 276

B


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

W

Cold finished

620

90

310

45

40(d)

40

...

P, Sh, St

Annealed

515

75

205

30

40

...

95 max

... A 580 A 167, A 240

Type 316L (UNS S31603) F

Annealed

450

65

170

25

40

50

...

A 473

B


480

70

170

25

40

50

...

A 276

B


620

90

310

45

30

40

...

A 276

B


480

70

170

25

30

40

...

A 276

W

Annealed

480

70

170

25

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

P, Sh, St

Annealed

485

70

170

25

40

...

95 max

A 167, A 240

Type 316Cb (UNS S31640) P, Sh, St

Annealed

515

75

205

30

30

...

95

A 167, A 240

B, Shapes


515

75

205

30

40

50

...

A 276

B, Shapes

Cold finished(b) and

620

90

310

45

30

40

...

A 276


Page 1326



01 Sep 2005

annealed B, Shapes


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

Type 317 (UNS S31700) B, F


515

75

205

30

40

50

...

A 276, A 473

B


620

90

310

45

30

40

...

A 276

B


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580 A 580

W

Cold finished

620

90

310

45

30(d)

40

...

P, Sh, St

Annealed

515

75

205

30

35

...

95 max

A 167, A 240

Type 317L (UNS S31703) B

Annealed

585(g)

85(g)

240(g)

35(g)

55(g)

65(g)

85 max(g)

P, Sh, St

Annealed

515

75

205

30

40

...

95 max

A 167

...

Type 317LM (UNS S31725) B, P

Annealed

515

75

205

30

40

...

...

A 276

P, Sh, St

Annealed

515

75

205

30

40

...

95 max

A 167

Types 321 (UNS S32100) and 321H (UNS 32109) B, F


515

75

205

30

40

50

...

A 276, A 473

B


620

90

310

45

30

40

...

A 276

B


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580

W

Cold finished

620

90

310

45

30(d)

40

...

A 580

P, Sh, St

Annealed

515

75

205

30

40

...

95 max

A 167, A 240

Types 347 (UNS S34700) and 348 (UNS S34800) B, F


515

75

205

30

40

50

B


620

90

310

45

30

40

...

A 276

A 276, A 473

B


515

75

205

30

30

40

...

A 276

W

Annealed

515

75

205

30

35(d)

50(d)

...

A 580 A 580

W

Cold finished

620

90

310

45

30(d)

40

...

P, Sh, St

Annealed

515

75

205

30

40

...

92 max

A 167, A 240

515

75

205

30

40

...

95 max

A 167, A 240

18-18-2 (UNS S38100) P, Sh, St

Annealed

Type 384 (UNS S38400) W(e)

Annealed

550

80

...

...

...

...

...

A 493

W(e)

Lightly drafted

585

85

...

...

...

...

...

A 493

35

30

50

...

B 473

20Cb-3 (UNS N08020), 20Mo-4 (UNS N08024), and 20Mo-6 (UNS N08026) B, W

Annealed

550

Shapes

Annealed

550

80

240

35

15

50

...

B 473

B, W

Annealed and strain hardened

620

90

415

60

15

40

...

B 473

W

Annealed and cold finished

620−830

90−120

...

...

...

...

...

B 473

P, Sh, St

Annealed

550

80

240

35

30

...

95 max

B 463

Pi, T

Annealed

550

80

240

35

30

...

...


80

240

B 464, B 468, B

Page 1327



01 Sep 2005

474, B 729 Sanicro 28 (UNS N08028) P, Sh, St

Annealed

500

73

215

31

40

...

Seamless Tube

Annealed

500

73

215

31

40

...

70−90(g) B 709 ...

B 668

...

B 511

Type 330 (UNS N08330) B

Annealed

485

70

210

30

30

...

P, Sh, St

Annealed

485

70

210

30

30

...

70−90(g) B 536

Pi

Annealed

485

70

210

30

30

...

70−90(g) B 535, B 546

AL-6X (UNS N08366) B, W

Annealed

515

75

210

30

30

...

...

B 691

P, Sh, St

Annealed

515

75

240

35

30

...

95 max

B 688

Pi, T

Annealed

515

75

210

30

30

...

...

B 675, B 676, B 690

Welded T

Cold worked

515

75

210

30

10

...

...

B 676

...

B 672

JS-700 (UNS N08700) B, W

Annealed

550

80

240

35

30

50

P, Sh, St

Annealed

550

80

240

35

30

...

75−90(g) B 599

Type 332 (UNS N08800) Pi, T

Annealed

515

75

210

30

30

...

...

B 163, B 407, B 514, B 515

Seamless Pi, T

Hot finished

450

65

170

25

30

...

...

B 407

B

Hot worked

550

80

240

35

25

...

...

B 408

B

Annealed

515

75

210

30

30

...

...

B 408

P

Hot rolled

550

80

240

35

25

...

...

B 409

P, Sh, St

Annealed

515

75

210

30

30

...

...

B 409

490

71

220

31

35

...

...

B 649

620−830

90−120

...

...

...

...

...

B 649

490

71

220

31

35

...

...

B 673, B 674, B 677

Type 904L (UNS N08904) B

Annealed

W

Cold finished

Pi, T

Annealed

P, Sh, St Annealed 490 71 220 31 35 ... 70−90(g) B 625 (a) B, bar; F, forgings; P, plate; Pi, pipe; Sh, sheet; St, strip; T, tube; W, wire. (b) Up to 13 mm (0.5 in.) thick. (c) Over 13 mm (0.5 in.) thick. (d) For wire 3.96 mm (5=32 in.) and under, elongation and reduction in area shall be 25 and 40%, respectively. (e) 4 mm (0.156 in.) in diameter and over. (f) For forged sections 127 mm (5 in.) and over, the tensile strength shall be 485 MPa (70 ksi). (g) For information only, not a basis for acceptance or rejection

Table 8 Minimum mechanical properties of high-nitrogen austenitic stainless steels Tensile strength Product form(a)

Condition

MPa

ksi

0.2% yield strength MPa

ksi

Elongation, Reduction % in area, %

Hardness HRB

ASTM specification

Type 201 (UNS S20100) B

Annealed

515

75

275

40

40

45

...

P, Sh, St

Annealed

655

95

310

45

40

...

100 max

Sh, St

1

860

125

515

75

25

...

...

A 666

Sh, St

1

1030

150

760

110

18

...

...

A 666

Sh, St

3

=4 hard

1210

175

930

135

12

...

...

A 666

Sh, St

Full hard

1280

185

965

140

9

...

...

A 666

B

Annealed

515

75

275

40

40

45

...

A 276

P, Sh, St

Annealed

620

90

260

38

40

...

...

A 666

Sh, St

1

860

125

515

75

12

...

...

A 660

=4 hard =2 hard

A 276 A 276, A 666

Type 202 (UNS S20200)

=4 hard


Page 1328



01 Sep 2005

Type 205 (UNS S20500) B, P, Sh, St

Annealed

790

115

450

65

40

...

10 max

A 666

690

100

380

55

35

55

...

A 276

Nitronic 50 (UNS S20910) B

Annealed

W

Annealed

690

100

380

55

35

55

...

A 580

Sh, St

Annealed

725

105

415

60

30

...

100 max

A 240

P

Annealed

690

100

380

55

35

...

100 max

A 240

725

105

380

55

40

...

...

A 666

Cryogenic Tenelon (UNS S21460) B, P, Sh, St

Annealed

Types 216 (UNS S21600) and 216L (UNS S21603) Sh, St

Annealed

690

100

415

60

40

...

100 max

A 240

P

Annealed

620

90

345

50

40

...

100 max

A 240

620

90

345

50

45

60

...

A 276, A 580

Nitronic 40 (UNS S21900) B, W

Annealed

21-6-9 LC (XM-11) (UNS S21904) B, W, shapes

Annealed

620

90

345

50

45

60

...

A 276, A 580

Sh, St

Annealed

690

100

415

60

40

...

...

A 666

P

Annealed

620

90

345

50

45

...

...

A 666 A 276, A 580


Annealed

690

100

380

55

30

50

...

Sh, St

Annealed

690

100

415

60

40

...

100 max

A 240

P

Annealed

690

100

380

55

40

...

100 max

A 240

690

100

380

55

30

50

...

A 276, A 580


Annealed

Type 304N (UNS S30451) B

Annealed

550

80

240

35

30

...

...

A 276

P, Sh, St

Annealed

550

80

240

35

30

...

92 max

A 240

Type 340HN (UNS S30452) B

Annealed

620

90

345

50

30

50

...

A 276

Sh, St

Annealed

620

90

345

50

30

...

100 max

A 240

P

Annealed

585

85

275

40

30

...

100 max

A 240 A 276

Type 304LN (UNS S30453) B

Annealed

515

75

205

30

...

...

...

P, Sh, St

Annealed

515

75

205

30

40

...

92 max

A 167, A 240

P, Sh, St

Annealed

600

87

310

45

40

...

95 max

A 167, A 240

B, shapes

Annealed

600

87

310

45

40

50

...

P, Sh, St

Annealed

650

94

300

44

35

...

96

B, shapes

Annealed

650

95

300

44

35

50

...

A 276

253 MA (UNS S30815) A 276

254 SMO (UNS S31254) A 167, A 240

Type 316N (UNS S31651) B

Annealed

550

80

240

35

30

...

...

A 276

P, Sh, St

Annealed

550

80

240

35

35

...

95 max

A 240

17-14-4 LN (UNS S31726) P, Sh, St

Annealed

550

80

240

35

40

...

96

A 167, A 240

B, shapes

Annealed

550

80

240

35

40

...

...

A 276

Annealed

550

80

240

35

40

...

95

A 167, A 240

317LN (UNS S31753) P, Sh, St AL 6XN (UNS N08367)


Page 1329



01 Sep 2005

B, W

Annealed

715

104

315

46

30

...

...

B691

P, Sh, St

Annealed

715

104

315

46

30

...

100

B688

Flanges, fittings, valves, Annealed and so on

715

104

315

46

30

50

...

B462

Seamless Pi, T

Annealed

715

104

315

46

30

...

...

B690

Welded Pi

Annealed

715

104

315

46

30

...

...

B676

Welded T

Solution treated and annealed

715

104

315

46

30

...

...

B676

Welded T

Cold worked

...

...

...

...

10

...

...

B676

Cronifer 1925 hMO (UNS N08925) B, W

Annealed

600

87

300

43

40

...

...

B649

Seamless Pi, T

Annealed

600

87

300

43

40

...

...

B677

Welded Pi

Annealed

600

87

300

43

40

...

...

B673

Welded T Annealed 600 87 (a) B, bar; P, plate; Pi, pipe; Sh, sheet; St, strip; T, tube; W, wire

300

43

40

...

...

B674

Certain austenitic stainless steels⎯the so-called metastable types⎯can develop higher strengths and hardnesses than other stable types for a given amount of cold work. In metastable austenitic stainless steels, deformation triggers and transformation of austenite to martensite. The effect of this transformation on strength is shown in Fig. 4 , which compares the stress-strain curve for stable type 304 with that for metastable type 301. The parabolic shape of the curve for type 304 indicates that strain hardening occurs throughout the duration of the application of stress, but that the amount of strain hardening for a given increment of stress decreases as stress increases. Fig. 4 Typical stress-strain curves for types 301 and 304 stainless steel

On the other hand, type 301 continues to strain harden well into the plastic range. The extended strain hardening is the result of the deformation-induced transformation of austenite to martensite. Ferritic types of stainless steel are defined as those that contain at least 10.5% Cr and that have microstructures of ferrite plus carbides. These steels are lower in toughness than the austenitic types. Basic room-temperature mechanical properties of ferritic stainless steels are given in Table 9 . Strength is enhanced only moderately by cold working. Additional specifications for ferritic stainless steels include: Product form

ASTM specification

Bar and wire

A 493

Billet and bar

A 314

Flanges, fittings and valves, and so on

A 182

Table 9 Minimum mechanical properties of ferritic stainless steels Tensile strength Product form(a)

Condition

0.2% yield strength

MPa

ksi

MPa

ksi

415

60

170

25

Elongation, Reduction % in area, %

Hardness, ASTM HRB specification

Type 405 (UNS S40500) B

Annealed


20

45

...

A 479

Page 1330



01 Sep 2005

F

Annealed

415

60

205

30

20

45

...

A 473

W

Annealed

480

70

280

40

20

45

...

A 580

P, Sh, St

Annealed

455

60

170

25

20

...

88 max

A 176, A 240

Type 409 (UNS S40900) P, Sh, St

Annealed

380

55

205

30

20

...

80 max

A 240

P, Sh, St

Annealed

380

55

205

30

22(c)

...

80 max

A 176

Type 429 (UNS S42900) B

Annealed

480

70

275

40

20

45

...

P, Sh, St

Annealed

450

65

205

30

22(c)

...

88 max

A 276 A 176, A 240

Type 430 (UNS S43000) B

Annealed

415

60

205

30

20

45

...

A 276

W

Annealed

480

70

275

40

20

45

...

A 580

P, Sh, St

Annealed

450

65

205

30

22(c)

...

88 max

485

70

275

40

20

45

...

A 473

...

...

...

...

...

A 581

A 176, A 240

Type 430F (UNS S43020) F

Annealed

W

Annealed

585−86 85−125 0

Type 439 (UNS S43035) B

Annealed

485

70

275

40

20

45

...

A 479

P, Sh, St

Annealed

450

65

205

30

22

...

88 max

A 240

515(b)

75(b)

310(b)

45(b)

30(b)

65(b)

...

...

Type 430Ti (UNS S43036) B

Annealed

Type 434 (UNS S43400) W

Annealed

545(b)

79(b)

415(b)

60(b)

33(b)

78(b)

90 max(b)

...

Sh

Annealed

530(b)

77(b)

365(b)

53(b)

23(b)

...

83 max(b)

...

530(b)

77(b)

365(b)

53(b)

23(b)

...

83 max(b)

... ...

Type 436 (UNS S43600) Sh, St

Annealed

Type 442 (UNS S44200) B

Annealed

550(b)

80(b)

310(b)

45(b)

20(b)

40(b)

90 max(b)

P, Sh, St

Annealed

515

75

275

40

20

...

95 max

A 176

415

60

275

40

20

...

95 max

A 176


Annealed

Type 446 (UNS S44600) B

Annealed, hot finished

480

70

275

40

20

45

...

A 276

B

Annealed, cold finished

480

70

275

40

16

45

...

A 276

W

Annealed

480

70

275

40

20

45

...

A 580

W


480

70

275

40

16

45

...

A 580

P, Sh, St

Annealed

515

75

275

40

20

...

95 max

A 176

Annealed

620(b)

90(b)

450(b)

65(b)

25(b)

...

90 min(b)

18 SR Sh, St

...

E-Brite 26-1 (UNS S44627) B


450

65

275

40

20

45

...

A 276

B


450

65

275

40

16

45

...

A 276

P, Sh, St

Annealed

450

65

275

40

22(c)

...

90 max


A 176, A

Page 1331



01 Sep 2005

240 MONIT (UNS S44635) P, Sh, St

Annealed

620

90

515

75

20

...

...

A 176, A 240

Annealed

585

85

450

65

18

...

100 max

A 176, A 240

Annealed

550

80

415

60

18

...

...

A 276, A 240

P, Sh, St

Annealed

550

80

415

60

20

...

98 max

A 176, A 240

B

Hot finished

480

70

380

55

20

40

...

A 276

B

Cold finished

520

75

415

60

15

30

...

A 276

Sea-Cure/SC-1 (UNS S44660) P, Sh, St 29-4C (UNS S44735) P, Sh, St 29-4-2 (UNS S44800)

B Annealed 480 70 380 55 20 40 ... A 479 (a) B, bar; F, forgings; W, wire; P, plate; Sh, sheet; St, strip. (b) Typical values. (c) 20% reduction for 1.3 mm (0.050 in.) and under in thickness

Duplex (Austenite/Ferrite) Types. Most wrought duplex stainless steels contain about 50% austenite-50% ferrite because of the balancing of elements that stabilize austenite (carbon, nitrogen, nickel, copper, and manganese) and ferrite (chromium, molybdenum, and silicon). Low carbon is maintained in most grades to minimize intergranular carbide precipitation. The austenite-ferrite balance provides wrought material with the optimum levels of mechanical properties and corrosion resistance. Because typically less austenite is present as-cast, welding consumables with enriched nickel are generally used to maintain austenite in the weld metal at levels generally similar to those in the base material. Yield strengths approximately twice that of type 316 can be obtained with annealed duplex stainless steels. Basic room-temperature mechanical properties of duplex stainless steel are given in Table 10 . Strength levels can be enhanced by cold working. Lower transverse ductility and impact strength can be expected because of the directional nature of the wrought micro-structure (typically elongated austenite islands in a ferrite matrix). Table 10 Minimum mechanical properties of duplex stainless steels Minimum values unless otherwise indicated Tensile strength Conditi Product form(a) on MPa ksi

0.2% yield strength MPa

ksi

ASTM Elongation, Reduction in Maximum specificat % area, % hardness, HRC ion

44LN (UNS S31200) F

Anneale d

690−900

100−130

450

65

25

50

...

A 182

P, Sh, St

Anneale d

690

100

450

65

25

...

220 HB

A 240

T

Anneale d

690

100

450

65

25

...

280 HB

A 789

Pi

Anneale d

690

100

450

65

25

...

280 HB

A 790

P, Sh, St

Anneale d

690

100

485

70

20

...

290 HB

A 240

T

Anneale d

690

100

450

65

25

...

30.5

A 789

Pi

Anneale d

690

100

450

65

25

...

...

A 790

T

Anneale d

630

92

440

64

30

...

30.5

A 789

Pi

Anneale d

630

92

440

64

30

...

30.5

A 790

DP-3 (UNS S31260)

3RE60 (UNS S31500)

2205 (UNS S31803)


Page 1332



01 Sep 2005

F

Anneale d

620

90

450

65

25

45

...

A 182

P, Sh, St

Anneale d

620

90

450

65

25

...

32

A 240

B, Shapes

Anneale d

620

90

448

65

25

...

290 HB

A 276

T

Anneale d

620

90

450

65

25

...

30.5

A 789

Pi

Anneale d

620

90

450

65

25

...

30.5

A 790

T

Anneale d

600

87

400

58

25

...

30.5

A 789

Pi

Anneale d

600

87

400

58

25

...

30.5

A 790

2304 (UNS S32304)

Ferralium 255 (UNS S32550) P, Sh, St

Anneale d

760

110

550

80

15

...

32

A 240

B, Shapes

Anneale d

760

110

550

80

15

...

297 HB

A 479

T

Anneale d

760

110

550

80

15

...

31.5

A 789

Pi

Anneale d

760

110

550

80

15

...

31.5

A 790


Anneale d

620

90

485

70

15

...

28

A 240

T

Anneale d

620

90

485

70

20

...

28

A 789

Pi

Anneale d

620

90

485

70

20

...

28

A 790

7-Mo PLUS (UNS S32950) P, Sh, St

Anneale d

690

100

480

70

15

...

31

A 240

B, Shapes

Anneale d

690

100

480

70

15

...

297 HB

A 479

T

Anneale d

690

100

480

70

20

...

30.5

A 789

Pi

Anneale 690 100 480 70 20 ... 30.5 A 790 d (a) B, bar; W, wire; P, plate; Sh, sheet; St, strip; T, tubing; Fl, flanges, fittings, valves, and parts for high-temperature service; Pi, pipe

Martensitic types are iron-chromium steels with or without small additions of other alloying elements. They are ferritic in the annealed condition, but are martensitic after rapid cooling in air or a liquid medium from above the critical temperature. Steels in this group usually contain no more than 14% Cr⎯except types 440A, 440B, and 440C, which contain 16 to 18% Cr⎯and an amount of carbon sufficent to permit hardening. If other elements are present, the total concentration is usually no more than 2 to 3%. Martensitic stainless steels may be hardened and tempered in the same manner as alloy steels. They have excellent strength and are magnetic. Basic room-temperature properties of the martensitic types are given in Table 11 and Fig. 5 . Additional specifications for martensitic stainless steels are: Product form

ASTM specification

Bar

A 582

Bolting

A 193

Nuts

A 194

Table 11 Minimum mechanical properties of martensitic stainless steels 0.2% yield


Reducti

Page 1333



Tensile strength Product form(a)

Condition

strength

MPa

ksi

MPa

ksi

485

70

275

40

01 Sep 2005

Elongat on ion, in area, Rockwell % % hardness

ASTM specification

Type 403 (UNS S40300) B, F


20

45

...

A 276, A 473, A 479

B


485

70

275

40

16

45

...

A 276

B

Intermediate temper, hot finished

690

100

550

80

15

45

...

A 276

B

Intermediate temper, cold finished

690

100

550

80

12

40

...

A 276

B

Hard temper, hot or cold finished

825

120

620

90

12

40

...

A 276

W

Annealed

485

70

275

40

20

45

...

A 580

W


485

70

275

40

16

45

...

A 580

W


690

100

550

80

12

40

...

A 580

W

Hard temper, cold finished

825

120

620

90

12

40

...

A 580

P, Sh, St

Annealed

485

70

205

30

25(b)

...

88 HRB max

A 176

485

70

275

40

20

45

...

Type 410 (UNS S41000) B, F


A 276, A 473, A 479

B


485

70

275

40

16

45

...

A 276

B


690

100

550

80

15

45

...

A 276

B


690

100

550

80

12

40

...

A 276

B

Hard temper, hot or cold finished

825

120

620

90

12

40

...

A 276

W

Annealed

485

70

275

40

20

45

...

A 580

W


485

70

275

40

16

45

...

A 580

W


690

100

550

80

12

40

...

A 580

W

Hard temper, cold finished

825

120

620

90

12

40

...

A 580

P, Sh, St

Annealed

450

65

205

30

22(b)

...

95 HRB max

A 176

P, Sh, St

Annealed

450

65

205

30

20

...

95 HRB max

A 240

Type 410S (UNS S41008) F

Annealed

450

65

240

35

22

45

...

P, Sh, St

Annealed

415

60

205

30

22(b)

...

88 HRB max

A 176, A 240

A 473

485

70

275

40

13

45

...

A 276, A 479

Type 410Cb (UNS S41040) B


B


485

70

275

40

12

35

...

A 276, A 479

B


860

125

690

100

13

45

...

A 276, 479

B


860

125

690

100

12

35

...

A 276, A 479

415

60

205

30

22

...

88 HRB max

A 276, A 240

E-4 (UNS S41050) P, Sh, St

Annealed

Type 414 (UNS S41400)


Page 1334



01 Sep 2005

B

Intermediate temper, cold or hot finished

795

115

620

90

15

45

...

A 276, A 479

W


1030 max

150 max

...

...

...

...

...

A 580

CA6NM (UNS S41500) P, Sh, St

Tempered

795

115

620

90

15

...

32 HRC max

A 176, A 240

B, F

Tempered

795

115

620

90

15

45

...

485

70

275

40

20

45

...

A 473

...

...

...

...

A 581

A 276, A 473, A 479

Types 416 (UNS S41600) and 416Se (UNS S41623) F

Annealed

W

Annealed

585−86 85−125 0

...

W

Intermediate temper

795−10 115−14 00 5

...

...

...

...

...

A 581

W

Hard temper

965−12 140−17 10 5

...

...

...

...

...

A 581

Type 416 plus X (UNS S41610) W

Annealed

585−86 85−125 0

...

...

...

...

...

A 581

W

Intermediate temper

795−10 115−14 00 5

...

...

...

...

...

A 581

W

Hard temper

965−12 140−17 10 5

...

...

...

...

...

A 581

760

110

15

45

...

A 565

Type 418 (UNS S41800) B, F

Tempered at 620 °C (1150 °F)

965

140

Type 420 (UNS S42000) B


1720

250

8(c)

25(c)

52 HRC(c)

W


860 max

125 max

...

...

...

...

...

A 580

P, Sh, St

Annealed

690

100

...

...

15

...

96 HRB max

A 176

1480(c) 215(c)

...

TrimRite (UNS S42010) W

Annealed

690 max

100 max

...

...

...

...

...

A 493

W

Lightly drafted

725 max

105 max

...

...

...

...

...

A 493

Type 422 (UNS S42200) B, F


825

120

585

85

17(d)

35

...

A 565

B, F


965

140

760

110

13

30

...

A 565

965

140

760

110

8

20

...

A 565

115

620

90

15

...

...

A 473

Lapelloy (UNS S42300) B, F


Type 431 (UNS S43100) F

Intermediate temper

795

F

Hard temper

1210

175

930

135

13

...

...

A 473

W


965 max

140 max

...

...

...

...

...

A 580

W

Annealed

760

110

...

...

...

...

...

A 493

W

Lightly drafted

795

115

...

...

...

...

...

A 493

105(c)

415(c)

Type 440A (UNS S44002) B

Annealed

725(c)

60(c)

20(c)

...

95 HRB(c)

...

B


1790(c) 260(c) 1650(c) 240(c)

5(c)

20(c)

51 HRC(c)

...

W


...

...

...

965


140

...

...

A 580

Page 1335



max

max 107(c)

01 Sep 2005

Type 440B (UNS S44003) B

Annealed

740(c)

62(c)

18(c)

...

96 HRB(c)

...

B


1930(c) 280(c) 1860(c) 270(c)

425(c)

3(c)

15(c)

55 HRC(c)

...

W


965 max

140 max

...

...

...

...

...

110(c)

450(c)

A 580

Type 440C (UNS S44004) B

Annealed

760(c)

65(c)

14(c)

...

97 HRB(c)

...

B


1970(c) 285(c) 1900(c) 275(c)

2(c)

10(c)

57 HRC(c)

...

W


965 140 ... ... ... ... ... A 580 max max (a) B, bar; F, forgings; P, plate; Sh, sheet; St, strip; W, wire. (b) 20% elongation for 1.3 mm (0.050 in.) and under in thickness. (c) Typical values. (d) Minimum elongation of 15% for forgings

Fig. 5 Typical hardnesses of selected martensitic stainless steels

Martensitic stainless steels harden when cooled off the mill after hot processing; therefore, they are often given a process anneal at 650 to 760 °C (1200 to 1400 °F) for about 4 h. Process annealing differs from full annealing, which is done by heating at 815 to 870 °C (1500 to 1600 °F), cooling in the furnace at a rate of 40 to 55 °C/h (75 to 100 °F/h) to about 540 °C (1000 °F), and then cooling in air to room temperature. Occasionally, martensitic types are purchased in the tempered condition; this condition is achieved by cooling directly off the mill to harden the steel and then reheating to a tempering temperature of 540 to 650 °C (1000 to 1200 °F), or by reheating the steel to a hardening temperature of 1010 to 1065 °C (1850 to 1950 °F), cooling it, and then tempering it. The influence of tempering temperature on the properties of hardened martensitic stainless steels is shown in Fig. 6 . In heat treating martensitic stainless steels, temperatures up to about 480 °C (900 °F) are referred to as stress-relieved temperatures because little change in tensile properties occurs upon heating hardened material to these temperatures. Temperatures of 540 to 650 °C (1000 to 1200 °F) are referred to as tempering temperatures, and temperatures of 650 to 760 °C (1200 to 1400 °F) are called annealing temperatures. Fig. 6 Variation of tensile properties and hardness with tempering temperature for three martensitic stainless steels


Page 1336



01 Sep 2005

Precipitation-hardening types are generally heat treated to final properties by the fabricator. Table 12 summarizes the minimum properties that can be expected in both material as received from the mill and material that has been properly heat treated. Table 12 Minimum mechanical properties of precipitation-hardening stainless steels Tensile strength Product form(a)

Condition

Yield strength

MPa

ksi

MPa

ksi

Elongat Reductio ion, n in % area, %

Hardness, HRC min

max

ASTM specification

PH 13-8 Mo (UNS S13800) B, F

H950

1520

220

1410

205

10

45;35(b)

45

...

A 564, A 705

B, F

H1000

1410

205

1310

190

10

50;40(b)

43

...

A 564, A 705

B, F

H1025

1275

185

1210

175

11

50;45(b)

41

...

A 564, A 705

B, F

H1050

1210

175

1140

165

12

50;45(b)

40

...

A 564, A 705

B, F

H1100

1030

150

930

135

14

50

34

...

A 564, A 705

B, F

H1150

930

135

620

90

14

50

30

...

A 564, A 705

B, F

H1150M

860

125

585

85

16

55

26

...

A 564, A 705

P, Sh, St

H950

1520

220

1410

205

6−10(c)

...

45

...

A 693

P, Sh, St

H1000

1380

200

1310

190

6−10(c)

...

43

...

A 693

15-5 PH (UNS S15500) B, F

H900

1310

190

1170

170

10;6(b) 35;15(b)

40

...

A 564, A 705

B, F

H925

1170

170

1070

155

10;7(b) 38;20(b)

38

...

A 564, A 705

B, F

H1025

1070

155

1000

145

12;8(b) 45;27(b)

35

...

A 564, A 705

B, F

H1075

1000

145

860

125

13;9(b) 45;28(b)

32

...

A 564, A 705

B, F

H1100

965

140

795

115

14;10(b 45;29(b)

31

...

A 564, A 705


Page 1337



01 Sep 2005

) B, F

H1150

930

135

725

105

16;11(b 50;30(b) )

28

...

A 564, A 705

B, F

H1150M

795

115

515

75

18;14(b 55;35(b) )

24

...

A 564, A 705

P, Sh, St

H900

1310

190

1170

170

5−10(c)

...

40

48

A 693

P, Sh, St

H1100

965

140

790

115

5−14(c)

...

29

40

A 693

1380

200

1210

175

7

25

...

...

A 564, A 705

PH 15-7 Mo (UNS S15700) B, F

RH950

B, F

TH1050

1240

180

1100

160

8

25

...

...

A 564, A 705

P, Sh, St

Annealed

1035 max

150 max

450 max

65 max

25 min

...

...

...

A 693

P, Sh, St

RH950(d)

1550

225

1380

200

1−4(c)

...

45−46

...

A 693

P, Sh, St

TH1050(d)

1310

190

1170

170

2−5(c)

...

40

...

A 693

P, Sh, St

Cold rolled condition C

1380

200

1210

175

1

...

41

...

A 693

P, Sh, St

CH900

1650

240

1590

230

1

...

46

...

A 693

40

...

A 564, A 705

17-4 PH (UNS S17400) B, F

H900(d)

1310

190

1170

170

10

40;35(e)

B, F

H925(d)

1170

170

1070

155

10

44;38(e)

38

...

A 564, A 705

B, F

H1025(d)

1070

155

1000

145

12

45

35

...

A 564, A 705

B, F

H1075(d)

1000

145

860

125

13

45

32

...

A 564, A 705

B, F

H1100(d)

965

140

795

115

14

45

31

...

A 564, A 705

B, F

H1150(d)

930

135

725

105

16

50

28

...

A 564, A 705

B, F

H1150M(d)

795

115

515

75

18

55

24

...

A 564, A 705

P, Sh, St

H900

1310

190

1170

170

5−10(c)

...

40

48

A 693

P, Sh, St

H1100

965

140

790

115

5−14(c)

...

29

40

A 693

1275

185

1030

150

6

10

41

...

A 564, A 705

17-7 PH (UNS S17700) B, F

RH950(d)

B, F

TH1050(d)

1170

170

965

140

6

25

38

...

A 564, A 705

P, Sh, St

RH950

1450(c)

210(c)

1310(c)

190(c)

1−6(c)

...

43(c)

44(c)

A 693

P, Sh, St

TH1050

1240(c)

180(c)

1030(c)

150(c)

3−7(c)

...

38

...

A 693

P, Sh, St

Cold rolled condition C

1380

200

1210

175

1

...

41

...

A 693

P, Sh, St

CH900

1650

240

1590

230

1

...

46

...

A 693

W

Cold drawn 1400−2035( 203−29 condition C c) 5(c)

...

...

...

...

...

...

A 313

W

CH900

...

...

...

...

...

...

A 313

1585−2515( 230−36 c) 5(c)

AM-350 (UNS S35000) P, Sh, St

Annealed

1380 max

200 max

585−620 max(c)

85−90 max(c)

8−12(c)

...

...

30

A 693

P, Sh, St P, Sh, St

H850

1275

185

1030

150

2−8(c)

...

42

...

A 693

H1000

1140

165

1000

145

2−8(c)

...

36

...

A 693

AM-355 (UNS S35500) F

H1000

1170

170

1070

155

12

25

37

...

A 705

P, Sh, St

H850

1310

190

1140

165

10

...

...

...

A 693

P, Sh, St

H1000

1170

170

1030

150

12

...

37

...

A 693

895(f)

130(f)

655

95

10

40

...

32

A 564(f)

Custom 450 (UNS S45000) B, shapes

Annealed


Page 1338


F, shapes

Annealed


860(f)

125(f)

655

95

B, F, shapes H900

1240(g)

180(g)

1170

170

B, F, shapes H950

1170(g)

170(g)

1100

B, F, shapes H1000

1100(g)

160(g)

B, F shapes H1025

1030(g)

B, F shapes H1050

1000(g)

B, F, shapes H1100 B, F, shapes H1150

10

01 Sep 2005

...

33

A 705(f)

6;10(b) 20;40(b)

39

...

A 564(g), A 705(g)

160

7;10(b) 22;40(b)

37

...

A 564(g), A 705(g)

1030

150

8;12(b) 27;45(b)

36

...

A 564(g), A 705

150(g)

965

140

34

...

A 564(g), A 705

145(g)

930

135

9;12(b) 30;45(b)

34

...

A 564(g), A 705

895(g)

130(g)

725

105

11;16(b 30;50(b) )

30

...

A 564(g), A 705

860(g)

125(g)

515

75

12−18( 35−55(h) h)

26

...

A 564(g), A 705

895−1205

130−16 5

620−1035

90−150

4 min

...

25

33

A 693

12

40

45

P, Sh, St

Annealed

P, Sh, St

H900

1240

180

1170

170

3−5(c)

...

40

...

A 693

P, Sh, St

H1000

1105

160

1035

150

5−7(c)

...

36

...

A 693

P, Sh, St

H1150

860

125

515

75

8−10(c)

...

26

...

A 693

B, F, shapes H900(i)

1620

235

1520

220

8

30

47

...

A 564(g), A 705(g)


1520

220

1410

205

10

40

44

...

A 564(g), A 705(g)


1410

205

1280

185

10

40

40

...

A 564(g), A 705(g)

Custom 455 (UNS S45500)

P, Sh, St H950 1530 222 1410 205 ... 44 ... A 693 ≤4 (a) B, bar; F, forgings; P, plate; Sh, sheet; St, strip; W, wire. (b) Higher value is longitudinal; lower value is transverse. (c) Values vary with thickness or diameter. (d) Longitudinal properties only. (e) Higher values are for sizes up to and including 75 mm (3 in.); lower values are for sizes over 75 mm (3 in.) up to and including 200 mm (8 in.) (f) Tensile strengths of 860 to 140 MPa (125 to 165 ksi) for sizes up to 13 mm (1=2in.). (g) Tensile strength only applicable up to sizes of 13 mm (1=2in.). (h) Varies with section size and test direction. (i) Up to and including 150 mm (6 in.)

The precipitation-hardening stainless steels are of two general classes: single-treatment alloys and double-treatment alloys. Single-treatment alloys, such as Custom 450, 17-4 PH, and 15-5 PH, are solution annealed at about 1040 °C (1900 °F) to dissolve the hardening agent. Upon cooling to room temperature, the structure transforms to martensite that is supersaturated with respect to the hardening agent. A single tempering treatment at about 480 to 620 °C (900 to 1150 °F) is all that is required to precipitate a secondary phase to strengthen the alloy. As listed in Table 12 , different tempering temperatures within this range produce different properties. Double-treatment alloys such as 17-7 PH are solution treated at about 1040 °C (1900 °F) and then water quenched to retain the hardening agent in solution in an austenitic structure. The austenite is conditioned by heating to 760 °C (1400 °F) to precipitate carbides and thereby unbalance the austenite so that it transforms to martensite upon cooling to a temperature below 15 °C (60 °F); this treatment produces condition T. Alternatively, the austenite may be conditioned at a higher temperature, 925 °C (1700 °F), at which fewer carbides precipitate, and then may be transformed to martensite by cooling to room temperature, followed by refrigerating to −75 °C (−100 °F); this treatment produces condition R. Transformation can also be effected by severe cold work (about 60 to 70% reduction); such treatment produces condition C. Once the structure has been transformed to martensite by one of these three processes, tempering at 480 to 620 °C (900 to 1150 °F) induces precipitation of a secondary metallic phase, which strengthens the alloy. Properties developed by typical TH, RH, and CH treatments are given in Table 12 for 17-7 PH.

Notch Toughness and Transition Temperature Notched-bar impact testing of stainless steels is likely to show a wide scatter in test results, regardless of type or test conditions. Because of this wide scatter, only general behavior of the different classes can be described. Austenitic types have good notched-bar impact resistance. Charpy impact energies of 135 J (100 ft · lb) or greater are typical of all types at room temperature. Cryogenic temperatures have little or no effect on notch toughness; ordinarily, austenitic stainless steel maintain values exceeding 135 J even at very low temperatures. On the other hand, cold work lowers the resistance to impact at all temperatures. Martensitic and ferritic stainless steels exhibit a decreasing resistance to impact with decreasing temperature, and the fracture appearance changes from a ductile mode at mildly elevated temperatures to a brittle mode at low temperatures. This fracture transition is characteristic of martensitic and ferritic materials. Both the upper-shelf energy and the lower-shelf energy are not


Page 1339



01 Sep 2005

greatly influenced by heat treatment in these stainless steels. However, the temperature range over which transition occurs is affected by heat treatment, minor variations in composition, and cold work. Heat treatments that result in high hardness move the transition range to higher temperatures, and those that result in low hardness move the transition range to lower temperatures. As indicated in Fig. 7 , transition generally occurs in the range of −75 to 95 °C (−100 to 200 °F), which is the temperature range in which martensitic stainless steels are ordinarily used. Consequently, it may be necessary to investigate fracture behavior thoroughly before specifying a martensitic or ferritic stainless steel for a particular application. Fig. 7 Typical transition behavior of type 410 martensitic stainless steel. All data from Charpy V-notch tests: A represents material tempered at 790 °C (1450 °F); final hardness, 95 HRB. B represents material tempered at 665 °C (1225 °F); final hardness, 24 HRC. C represents material tempered at 595 °C (1100 °F); final hardness, 30 HRC.

Fracture toughness data are not available for many of the standard types of stainless steel. Most testing has been concentrated on the high-strength precipitation-hardening stainless steels because these materials have been used in critical applications where fracture toughness testing has been found most useful for evaluating materials. Table 13 lists typical fracture toughness for several of the high-strength stainless steels for which this property has been determined. Table 13 Longitudinal fracture toughness of precipitation-hardening (PH) stainless steels Designation

Condition(a)

Hardness, HRC

Fracture toughness p p MPa m ksi in:

17-4 PH

H900

44

53

48

17-7 PH

RH950

44

76

69

Custom 450

Aged at 480 °C (900 °F)

43

81

74

Custom 455

Aged at 480 °C (900 °F)

50

47

43

Aged at 510 °C (950 °F)

48

80

73

Aged at 540 °C (1000 °F)

44

110

100

H950

47

99

90

46

121

110

42

55

50

PH 13-8 Mo

H1000 PH 15-7 Mo TH1080 (a) H, hardened; RH, refrigeration hardened; TH, transformation hardened

Fatigue Strength Three types of fatigue tests are used to develop data on the fatigue behavior of stainless steels: • Rotating-beam test, the most commonly used, which most closely approximates the kind of loading to which shafts and axles are subjected • Flexural fatigue test (used to evaluate the behavior of sheet), which most closely simulates the action of leaf springs, which are expected to flex without deforming or breaking • Axial-load fatigue test, which subjects a fatigue specimen to unidirectional loading that can range from full reversal (tension-compression) to tension-tension loading and can have virtually any conceivable ratio of maximum stress to minimum


Page 1340



01 Sep 2005

stress Fatigue data can be given in the form of stress-number of cycles (S-N) curves (Fig. 8 ) or constant-life diagrams (Fig. 9 ). Data from any of the three types of tests can be presented as S-N curves, but only data from flexural fatigue and axial fatigue tests can be presented in the form of a constant-life diagram. In analyzing fatigue data, and particularly in selecting materials on the basis of fatigue life, it is important to understand the influence of the stress ratio on fatigue life. In general, fatigue conditions involving tension-compression loading (stress ratio, R, between 0 and −1) lead to shorter fatigue lives than conditions involving tension-tension loading (stress ratio, R, between 0 and +1) at the same value of maximum stress. Fig. 8 Typical rotating-beam fatigue behavior of types 304 and 310 stainless steel

Fig. 9 Constant-life fatigue diagram for PH 13-8 Mo stainless steel, condition H1000

Elevated-Temperature Properties Many stainless steels⎯particularly the austenitic types 304, 309, 310, 316, 321, and 347 and certain precipitation-hardening types such as PH 15-7 Mo, 15-5 PH, 17-4 PH, 17-7 PH, AM-350, and AM-355⎯are used extensively for elevated-temperature applications such as chemical processing equipment, high-temperature heat exchangers, and superheater tubes for power boilers. For more details on the elevated-temperature properties of selected types, see the article "Elevated-Temperature Properties of Stainless Steels" in this Volume. Extended service at elevated temperature can result in the embrittlement of stainless steels or in sensitization, which degrades the ability of the material to withstand corrosion, particularly in acid media. Most often, such degradation is caused by the precipitation of secondary phases such as carbides, α′ phase, or σ phase. Precipitation depends on both time and temperature; longer times at temperature and higher temperatures within the precipitation temperature range promote more extensive precipitation. The problems arising from embrittlement and sensitization and the remedies that can help combat them are discussed in detail in the section "Fabrication Characteristics" in this article.

Subzero-Temperature Properties Austenitic stainless steels have been used extensively for subzero applications to −270 °C (−450 °F). These steels contain sufficient amounts of nickel and manganese to depress the temperature at which martensite starts to form from austenite upon cooling, Ms, into the subzero range. Thus they retain face-centered cubic (fcc) crystal structures upon cooling from hot working or annealing temperatures. Yield and tensile strengths of austenitic stainless steels increase substantially as testing temperature is decreased, and these steels retain good ductility and toughness at −270 °C (−450 °F). Most austenitic stainless steels may be readily fabricated by welding, but sometimes the welding heat causes sensitization that reduces corrosion resistance in the weld


Page 1341



01 Sep 2005

area. The strength of austenitic steels can be increased by cold rolling or cold drawing. Cold working at −195 °C (−320 °F) is more effective in increasing strength than cold working at room temperature. For metallurgically unstable stainless steels such as 301, 304, and 304L, plastic deformation at subzero temperatures causes partial transformation to martensite, which increases strength. For some cryogenic applications, it is desirable to use a stable stainless steel such as type 310. Compositions of austenitic stainless steels of interest are presented in Table 3 . Small amounts of nitrogen increase the strengths of these steels. Manganese additions are used in some steels to replace some of the nickel. Type 416 is a martensitic chromium stainless steel that is usually used in the quenched and tempered condition. It is included in this series because there are applications in rotating pumps and other machinery in which a magnetic material is needed to activate counters. Types 301 and 310 have been used in the form of extrahard cold-rolled sheet to provide high strength in such applications as the liquid oxygen and liquid hydrogen tanks for Atlas and Centaur rockets. Joining was done by butt fusion welding, and reinforcing strips were spot welded to the tank along the weld joint. In another method for producing high-strength cylindrical tanks, welded preform tanks are fabricated from annealed type 301 stainless steel, submerged in liquid nitrogen while in a cylindrical die, and expanded (cryoformed) by pressurizing until the preform fits the die. The amount of strengthening depends on the amount of plastic deformation incurred in expanding the preform to the size of the die. Strengthening results from the dual effects of the cold working of the austenite and the partial transformation of the austenite to martensite. Type 304 stainless steel is usually used in the annealed condition for tubing, pipes, and valves employed in the transfer of cryogens; for Dewar flasks and storage tanks; and for structural components that do not require high strength. Types 310 and 310S are considered metallurgically stable for all conditions of cryogenic exposure. Therefore, these steels are used for structural components in which maximum stability and a high degree of toughness are required at cryogenic temperatures. Type 316 stainless steel is less stable than type 310, but tensile specimens of type 316 pulled to 0.2% offset (at the yield load) at −270 °C (−450 °F) showed no indication of martensite formation in the deformed regions (Ref 4). However, when tensile specimens of type 316 were pulled to fracture at −270 °C (−450 °F), the metallographic structures in the areas of the fractures transformed to approximately 50% martensite (Ref 5). Type 316 stainless steel is an important candidate material for structural components of superconducting and magnetic fusion machinery. For higher-strength components of cryogenic structures, there are several stainless steels that contain significant amounts of manganese in place of some of the nickel, along with small additions of nitrogen and other elements that increase strength. Among these stainless steels are 21-6-9, Pyromet 538, Nitronic 40, and Nitronic 60. Tensile Properties. Typical tensile properties of annealed 300 series austenitic stainless steels at room temperature and at subzero temperatures are presented in Table 14 , and tensile properties of cold worked 300 series stainless steels are given in Table 15 . Cold working substantially increases yield and tensile strengths and reduces ductility, but ductility and notch toughness of the cold-worked alloy are often sufficient for cryogenic applications. Tensile properties of other stainless steels are presented in Table 16 . For the annealed alloys, the greatest effect of the nitrogen addition is to produce an increase in yield strength at cryogenic temperatures. The data for cold-worked AISI 202, a nitrogen-strengthened stainless steel, indicates how this alloy can be strengthened by cold working that results in reduced ductility. Because of its low ductility, alloy 416 is not recommended for use below −196 °C (−320 °F) except in nonstressed applications. Table 14 Typical tensile properties of annealed type 300 austenitic stainless steels Tensile strength

Temperature °C

°F

MPa

Yield strength

ksi

MPa

ksi

Elongatio n, Reduction, % in area, %

Notch tensile strength(a)

Young's modulus

MPa

ksi

GPa

106 psi

303 bar, longitudinal orientation 24

75

730

106

425

61.4

67

70

...

...

...

...

−78

−108

1190

172

435

63.3

43

60

...

...

...

...

−196

−320

1660

240

465

67.3

36

54

...

...

...

...

−253

−423

2060

298

570

82.6

33

...

...

...

...

...

−269

−452

1830

266

...

...

30

37

...

...

...

...

304 sheet, longitudinal orientation 24

75

660

95.5

295

42.5

75

...

715

104

...

...

−196

−320

1625

236

380

55.0

42

...

1450

210

...

...

−253

−423

1800

261

425

62.0

31

...

1160

168

...

...

−269

−452

1700

247

570

82.5

30

...

1230

178

...

...

304 plate, longitudinal orientation 24 −253

75

590

85.9

330

47.6

64

...

...

...

...

...

−423

1720

250

410

59.4

...

...

...

...

...

...

304 bar, longitudinal orientation


Page 1342



01 Sep 2005

24

75

640

92.8

235

33.9

76

82

710

103

...

...

−78

−108

1150

167

300

43.2

50

76

...

...

...

...

−196

−320

1520

221

280

40.9

45

66

1060

153

...

...

−253

−423

1860

270

420

60.6

27

54

1120

162

...

...

−269

−452

1720

250

400

58.2

30

55

...

...

...

...

304L sheet, longitudinal orientation 24

75

660

95.9

295

42.8

56

...

730

106

...

...

−78

−108

980

142

250

36.0

43

...

1030

150

...

...

−196

−320

1460

212

275

39.6

37

...

1420

206

...

...

−253

−423

1750

254

305

44.5

33

...

1290

187

...

...

−269

−452

1590

230

405

58.5

29

...

1460

212

...

...

223

410

59.5

35

...

...

...

...

...

190

27.6

304L sheet, transverse orientation −269

−452

1540

304L bar, longitudinal orientation 24

75

660

95.5

405

58.9

78

81

...

...

−78

−108

1060

153

435

62.8

70

74

...

...

...

...

−196

−320

1510

219

460

66.6

43

66

...

...

205

29.7

−253

−423

1880

273

525

75.8

42

41

...

...

...

...

−269

−452

1660

241

545

79.4

34

56

...

...

200

29.2


75

570

83.0

240

35.0

50

...

645

93.9

...

...

−196

−320

1080

156

545

79.1

68

...

1070

155

...

...

−253

−423

1300

188

715

104

56

...

1250

182

...

...

−269

−452

1230

178

770

112

58

...

...

...

...

...

75

600

86.8

240

34.8

46

...

630

91.6

...

...

−452

1280

186

800

116

58

...

...

...

...

...

310 sheet, transverse orientation 24 −269


75

585

84.8

340

49.1

50

76

770

112

...

...

−78

−108

740

107

305

43.9

72

68

...

...

...

...

−196

−320

1090

158

520

75.5

68

50

...

...

205

29.9

−253

−423

1390

202

855

124

44

48

1305

189

...

...

−269

−452

1300

189

715

104

50

41

...

...

205

29.9

310S forging, transverse orientation 24

75

585

84.8

260

37.9

54

71

800

116

...

...

−196

−320

1100

159

605

87.6

72

52

1350

196

...

...

−269

−452

1300

189

815

118

64

45

1600

232

...

...

316 sheet, longitudinal orientation 24 −253

75

595

86.4

275

39.8

60

...

...

...

...

...

−423

1580

229

664

96.6

55

...

...

...

...

...


75

620

89.6

225

32.4

55

...

625

90.4

180

26.0

−196

−320

1380

200

315

45.6

46

...

1520

220

205

29.5

−253

−423

1650

239

375

54.5

36

...

1460

212

210

30.7


75

675

97.6

430

62.2

55

79

...

...

...

...

−78

−108

1060

153

385

55.9

46

73

...

...

...

...


Page 1343



01 Sep 2005

−196

−320

1540

223

450

65.4

38

60

...

...

...

...

−253

−423

1860

270

405

58.5

35

44

...

...

...

...


75

650

94

255

37

52

...

...

...

...

...

−196

−320

1365

198

420

61

47

...

...

...

...

...

−253

−423

1610

234

435

63

35

...

...

...

...

...

24

75

670

97.4

340

49.3

57

76

...

...

...

...

−78

−108

995

144

475

68.8

51

71

...

...

...

...

−196

−320

1470

214

430

62.2

43

60

...

...

...

...

347 bar

1850 268 525 76.4 38 45 ... ... ... ... −253 −423 (a) Stress concentration factor, Kt, is 5.2 for 304 and 304L sheet, 14 for 304 bar, 6.3 for 310 sheet, 6.4 for 310 bar; Kt is 10 for 310S forging; Kt is 3.5 for 321 sheet. Source: Ref 6, 7, 8, 9, 10, 11, 12

Table 15 Typical tensile properties of cold-worked type 300 austenitic stainless steel sheet Temperature °C

°F

Tensile strength

Yield strength

MPa

MPa

ksi

ksi


Young's modulus

Elongation, %

MPa

ksi

GPa

106 psi

301, hard, cold rolled (42−60% reduction), longitudinal orientation 24

75

1310

190

1200

174

18

1390

201

...

...

−78

−108

1560

226

1130

164

23

1460

212

...

...

−196

−320

2020

293

1380

200

19

1660

241

...

...

−253

−423

2110

306

1610

233

14

1830

265

...

...

301, hard, cold rolled (42−60% reduction), transverse orientation 24

75

1310

190

1060

153

10

1430

207

...

...

−78

−108

1560

226

1070

155

28

1430

208

...

...

−196

−320

2060

299

1310

190

28

1670

243

...

...

−253

−423

1900

275

1570

227

8

1360

197

...

...

301, extra hard, cold rolled (>60% reduction), longitudinal orientation 24

75

1500

217

1370

198

9

1600

232

175

25.6

−78

−108

1710

248

1400

203

22

1680

244

180

26.3

−196

−320

2220

322

1610

234

22

1940

282

180

26.2

−253

−423

2220

322

1810

262

13

1890

274

190

27.6

−269

−452

2140

310

1930

280

2

...

...

...

...

301, extra hard, cold rolled (>60% reduction), transverse orientation 24

75

1590

230

1280

186

8

1520

220

...

...

−78

−108

1770

257

1250

181

18

1590

230

...

...

−196

−320

2190

318

1560

226

18

1680

244

...

...

−253

−423

2180

316

1830

266

5

1340

194

...

...

304, hard, cold rolled, longitudinal orientation 24

75

1320

191

1190

173

3

1460

212

180

25.9

−78

−108

1470

213

1300

188

10

1590

231

185

26.9

−196

−320

1900

276

1430

208

29

1910

277

200

29.1

−253

−423

2010

292

1560

226

2

2160

313

210

30.5

304, hard, cold rolled, transverse orientation 24

75

1440

209

1180

171

5

1200

174

195

28.0

−78

−108

1600

232

1330

193

7

1400

203

200

28.9

−196

−320

1870

271

1480

214

23

1690

245

205

30.0


Page 1344


−253

−423


2160

313

1560

226

01 Sep 2005

1

1900

276

215

31.1

304L, 70% cold reduced, longitudinal orientation 24

75

1320

192

1080

156

3

...

...

...

...

−196

−320

1770

256

1530

222

14

...

...

...

...

−253

−423

1990

288

1770

256

2

...

...

...

...

304L, 70% cold reduced, transverse orientation 24

75

1440

209

1220

177

4

...

...

...

...

−196

−320

1890

274

1630

236

12

...

...

...

...

−253

−423

2230

324

1940

282

1

...

...

...

...

3

1360

197

175

25.4

310, 75% cold reduced, longitudinal orientation 24

75

1180

171

1100

160

−78

−108

1410

204

1290

187

4

1530

222

175

25.5

−196

−320

1720

249

1540

223

10

1900

276

180

26.4

−253

−423

2000

290

1790

259

10

2230

324

195

28.3

310, 75% cold reduced, transverse orientation 24

75

1370

199

1110

161

4

1370

199

195

28.1

−78

−108

1540

224

1290

187

8

1640

238

190

27.6

−196

−320

1880

272

1520

221

10

2050

297

195

28.2

2140 −253 −423 (a) Kt = 6.3. Source: Ref 6, 10, 13

311

1790

260

9

2190

318

200

29.1

Table 16 Typical tensile properties of stainless steels other than type 300 series steels Temperature °C

°F


ksi

Yield strength MPa

ksi

Elongation, Reduction, % in area, %


Young's modulus

MPa

ksi

GPa

106 psi

202 sheet, annealed, longitudinal orientation 24

75

705

102

325

47.1

57

...

...

...

...

...

−73

−100

1080

156

485

70.2

41

...

...

...

...

...

−196

−320

1590

231

610

88.3

52

...

...

...

...

...

−268

−450

1420

206

765

111

25

...

...

...

...

...

202 sheet, cold reduced 50%, longitudinal orientation 24

75

1080

156

965

140

21

...

...

...

...

...

−196

−320

1970

286

1070

155

28

...

...

...

...

...

−268

−450

1950

283

1240

180

20

...

...

...

...

...

705

102

385

55.9

54

80

...

...

...

...

21-6-9 plate, longitudinal orientation(b) 24

75

−78

−108

895

130

590

85.4

60

75

...

...

...

...

−196

−320

1510

219

970

141

41

33

...

...

...

...

−253

−423

1660

241

1220

177

16

26

...

...

...

...

−269

−452

1700

247

1350

196

22

30

...

...

...

...

Pyromet 538 plate, longitudinal orientation(c) 24

75

675

97.9

340

49.0

75

81

...

...

...

...

−196

−320

1370

199

800

116

76

73

...

...

...

...

−269

−452

1490

216

1010

147

52

59

...

...

...

...

Nitronic 40 plate, electroslag remelted; as-rolled 24

75

1010

146

840

122

35

72

...

...

...

...

−73

−100

1170

169

945

137

36

71

...

...

...

...

−196

−320

1830

266

1540

223

31

64

...

...

...

...


Page 1345



01 Sep 2005

Nitronic 60 bar, annealed, longitudinal orientation 24

75

750

109

400

58.1

66

79

1080

157

165

24.0

−73

−100

1020

148

535

77.9

70

81

1480

215

165

24.2

−196

−320

1500

218

695

101

60

66

1900

275

170

24.8

−253

−423

1410

204

860

125

24

27

1870

271

170

24.8

416 bar, longitudinal orientation(d) 24

75

1400

203

1200

174

15

53

...

...

...

...

−78

−108

1500

218

1260

183

15

52

...

...

...

...

−196

−320

1800

261

1600

232

9

24

...

...

...

...

2020 293 2020 293 0.4 2 ... ... ... ... −253 −423 (a) Kt = 7 for Nitronic 60 bar. (b) Annealed 1 h at 1065 °C (1950 °F), water quenched. (c) Annealed 1 h at 1095 °C (2000 °F), water quenched. (d) Heat treatment: 1 h at 980 °C (1800 °F), oil quenched, tempered 4 h at 370 °C (700 °F), air cooled. Source: Ref 6, 10, 11, 14, 15, 16, 17, 18, 19, 20

Results of tensile tests on stainless steel weldments at subzero temperatures, given in Table 17 , may be significant in selecting stainless steels for cryogenic applications. Results of ultrasonic determinations of Young's modulus and Poisson's ratio for three stainless steels, shown in Fig. 10 and 11 , serve to supplement the tensile data. Table 17 Typical tensile properties of stainless steel weldments

Alloy condition Type 301, cold rolled 60%; tested as welded

3

Type 310, =4 hard; tested as welded

AISI, 310S, annealed

21-6-9, annealed

Weld ing proce ss Filler GTA

GTA

SMA

Test Yield Base metal temperature strength orientation °C °F MPa ksi Form (a)

None Sheet

310

310S

L

Sheet

L

Plate

...


ksi

Notch Reduct tensile Elonga ion strength(b) in area, tion, % % MPa ksi

24

75

...

...

1034

150

7

...

...

...

−78

−108

...

...

1489

216

13

...

...

...

−196 −320

...

...

2006

291

16

...

...

...

−253 −423

...

...

1675

243

6

...

...

...

24

75

380

55.1

530

76.8

4

...

...

...

−78

−108

523

75.9

723

105

4

...

...

...

−196 −320

752

109

1026

149

4

...

...

...

334

48.5

582

84.4

40

76

841

122

−196 −320

24

75

660

96.6

1066

155

46

67

1428

207

−269 −452

829

120

1102

160

26

24

1672

242

SMA Inconel Plate 625

Weld(c)

−269 −452

878

127

1276

185

31

27

...

...

HAZ(c)

−269 −452 1728

251

1873

272

21

33

...

...

GTA Inconel Plate 625

Weld(c)

−269 −452

951

138

1222

177

18

20

...

...

HAZ(c)

−269 −452 1740

252

1921

279

17

37

...

...

GMA Inconel Plate 625

Weld(c)

−269 −452

121

1087

158

19

27

...

...

HAZ(c)

−269 −452 1689

GTA Pyrome Plate t 538

...

GMA IN-182 Plate

...

833

245

1866

271

15

27

414

60.0

725

105

51

74

1238

180

−196 −320 1009

146

1456

211

48

61

2119

307

−269 −452 1240

180

1646

239

31

24

1841

267

24

24

75

75

−196 −320

413

59.9

729

106

53

75

1018

148

800

116

1045

152

6

37

1416

205

6 40 1419 206 −269 −452 805 117 1086 158 (a) L, longitudinal. (b) Kt = 10. (c) Weld parallel with specimen axis; weld specimens were all weld metal; HAZ specimens contained HAZ plus some weld metal and some base metal. Source: Ref 6, 15, 17, 18, 21, 22

Fig. 10 Young's modulus for three austenitic stainless steels as determined ultrasonically. Source: Ref 23


Page 1346



01 Sep 2005

Fig. 11 Poisson's ratios for three austenitic stainless steels as determined ultrasonically. Source: Ref 23

Fracture Toughness. Fracture toughness data for stainless steels are limited because steels of this type that are suitable for use at cryogenic temperatures have very high toughness. The fracture toughness data that are available were obtained by the J-integral method and converted to KIc(J) values. Such data for base metal and weldments are shown in Table 18 . Fracture toughness of base metals are relatively high even at −269 °C (−452 °F); fracture toughness of fusion zones (FZ) of welds may be lower or higher than that of the base metal. Table 18 Fracture toughness of austenitic stainless steels and weldments for compact tension specimens

Alloy and condition(a) Type 310S, annealed

Form Plate

Room-temperat ure yield strength Orientatio n MPa ksi

Fracture toughness, (KIc), J, at 24 °C (75 °F) p p MPa m ksi in:

−196 °C (−320 °F) p p MPa m ksi in:

−269 °C (−452 °F) p p MPa m ksi in:

261

37.9

T-L

...

...

...

...

262

236

...

...

...

...

...

...

...

118

106

338

49

T-L

...

...

275

250

182

165

...

...

...

...

...

...

...

82.4

74.4

Weldmen ... ... ... ... ... ... ... 176 t (a) STQ, solution treated and quenched. Filler wires for 310S: E 310-16; For Pyromet 538: 21-6-9. Source: Ref 17, 24, 25, 26

159

Weldmen t Pyromet 538, STQ

Plate Weldmen t

Fracture Crack Growth Rates. Available data for determining fatigue crack growth rates at room temperature and at subzero temperatures for austenitic stainless steels and weldments are presented in Table 19 . The fatigue crack growth rates of the base metals are generally higher at room temperature than at subzero temperatures, or about equal at room temperature and at subzero temperatures, except for 21-6-9 stainless steel. For 21-6-9, fatigue crack growth rates are higher at −269 °C (−452 °F) than at room temperature. A log-log plot of the da/dN data for type 304 stainless steel is shown in Fig. 12 . For this steel, fatigue crack growth rates are nearly the same, at the same values of ∆K, for room-temperature and cryogenic-temperature tests. Fatigue crack growth rates in the fusion zones of welds tend to be higher than in the base metal. Table 19 Fatigue crack growth rate (da/dN) data for compact tension specimens of austenitic stainless steels Test temperature or temperature range


C(b)

Estimated range for ∆K

Page 1347


Alloy and condition


Orien ta Fre Stress tion(a quency, ratio, ) Hz R

°C

°F

01 Sep 2005

da/dN:mm/ da/dN:in./ cycle cycle ∆K:MPa p p m ∆K:ksi in: n(b)

p MPa m

p ksi in:

Type 304 annealed plate

T-L

20−28

0.1

24 to −269

75 to −452

2.7 × 10−9 1.4 × 10−10

3.0

22−80

20−73

Type 304L annealed plate

T-L

20−28

0.1

24

75

2.0 × 10−10 1.2 × 10−11

4.0

22−54

20−49

−196, −269 −320, −452 3.4 × 10−11 2.0 × 10−12

4.0

26−80

24−73

Type 310S annealed plate

T-L

20−28

0.1

24

75

3.5 × 10−11 2.1 × 10−12

4.4

24−35

22−32

24

75

4.7 × 10−9 2.4 × 10−10

3.0

35−60

32−55

−196, −269 −320, −452 1.1 × 10−10 6.1 × 10−12

3.7

25−80

23−73

...

10

0.1

−196, −269 −320, −452 1.4 × 10−10 7.9 × 10−12 3.75

24−71

22−65

...

10

0.1

−196, −269 −320, −452 7.8 × 10−13 5.0 × 10−14 5.15

27−66

25−60

Type 316 annealed plate

T-L

20−28

0.1

24 to −269

3.8

19-16

17-14

21-6-9 annealed plate

T-L

20−28

0.1

24, −196

75, −320

1.9 × 10−10 1.1 × 10−11

3.7

25−80

23−73

−269

−452

3.6 × 10−11 2.2 × 10−12

4.4

25−70

23−64

24

75

1.8 × 10−10 9.9 × 10−12

3.7

26−55

24−50

6.36

24−44

22−40

Type 310S, SMA weld with E310-16 filler

Pyromet 538, GTA weld in annealed plate using 21-6-9 filler

T-L

10

0.1

75 to −452 2.1 × 10−10 1.2 × 10−11

−196, −269 −320, −452 7.6 × 10−14

5.47 × 10−15

T-L 10 0.1 24 to −269 75 to −452 2.5 × 10−12 1.6 × 10−13 5.13 Pyromet 538, SMA weld in 25−55 23−50 annealed plate using Inconel 182 filler (a) T, transverse; L, longitudinal. (b) C and n are constants from da/dN = C (∆K)″; ∆K, stress intensity factor range. Source: Ref 17, 26

Fig. 12 Fatigue crack growth rate data for type 304 austenitic stainless steel (annealed) at room temperature and at subzero temperatures. Source: Ref 27

Fatigue Strength. The results of flexural and axial fatigue tests at 106 cycles on austenitic stainless steels at room temperature and at subzero temperatures are presented in Table 20 . Fatigue strength increases as exposure temperature is decreased. Notched specimens have substantially lower fatigue strengths than corresponding unnotched specimens at all testing temperatures. Reducing the surface roughness of unnotched specimens improves fatigue strength.


Page 1348



01 Sep 2005

Table 20 Results of fatigue life tests on austenitic stainless steels Fatigue strengths at 106 cycles

Alloy and condition

Stressin g Stress mode ratio, R

Cyclic frequency, Hz

Type 301 sheet, extra full hard

Flex

−1.0

29, 86

Type 304L bar, annealed

Axial

−1.0

...

24 °C (75 °F)

−196 °C (−320 °F)

−253 °C (−423 °F)

Kt

MPa

ksi

MPa

ksi

MPa

ksi

1

496

72

793

115

669

97

3.1

172

25

303

44

...

...

1

269

39

483

70

552(a)

80(a)

3.1

193

28

207

30

228(a)

33(a)

Flex(b)

−1.0

...

1

186

27

455

66

597

84

Flex(c)

−1.0

...

1

213

31

490

71

662

96

Type 310 bar, annealed

Axial

−1.0

...

1

255

37

469

68

607(a)

88(a)

3.1

186

27

234

34

352(a)

51(a)

Type 321 sheet, annealed

Axial

−1.0

...

1

221

32

303

44

372

54



Flex(b)

−1.0

Flex(b)

−1.0

3.5

124

18

154

22.3

181

26.3

30−40

1

172

25

303

44

358

52

30−40

1

221

32

421

61

386

56

68

510

74

Flex(c) 1 241 35 469 −1.0 30−40 (a) Tested at −269 °C (−452 °F). (b) Surface finish 64 rms. (c) Surface finish 11 rms. Source: Ref 6, 28,29, 30, 31

Influence of Product Form on Properties The mechanical properties of cast or wrought stainless steels vary widely from group to group, vary less widely from type to type within groups, and may vary with product form for a given type. Because of the wide variation from group to group, one must first decide whether a martensitic, ferritic, austenitic, duplex, or precipitation-hardening stainless steel is most suitable for a given application. Once the appropriate group is selected, the method of fabrication or service conditions may then dictate which specific type is required. Before typical properties of the various product forms are discussed, it is important that two key points about stainless steels be recognized. First, many stainless steels are manufactured and/or used in a heat-treated condition, that is, in some thermally treated condition other than process annealed or, typically, mill processed. When this is the case, a tabulation of typical properties may not give all the required information. Second, in many products strain hardening during fabrication is a very important consideration. All stainless steels strain harden to some degree depending on structure, alloy content, and amount of cold working. Consequently, for applications in which the service performance of the finished product depends on the enhancement of properties during fabrication, it is essential that the manufacturer determine this effect independently for each individual product. Here, techniques such as statistical-reliability testing are invaluable. Cast Structures. Whether produced as ingot, slab, or billet in a mill or as shape castings in a foundry, cast structures can exhibit wide variations in properties. Because of the possible existence of large dendritic grains, intergranular phases, and alloy segregation, typical mechanical properties cannot be stated precisely and generally are inferior to those of any wrought structure. Detailed information on the composition and properties of cast stainless steels is given in the article "Cast Stainless Steels" in this Volume. Hot Processing. The initial purpose of hot rolling or forging an ingot, slab, or billet is to refine the cast structure and improve mechanical properties. Hot reduced products and hot reduced and annealed products exhibit coarser grain structures and lower strengths than cold processed products. Grain size and shape depend chiefly on start and finish temperatures and on the method of hot reduction. For instance, cross-rolled hand mill plate will exhibit a more equiaxed grain structure than continuous hot rolled strip. Hot reduction may be a final sizing operation, as in the case of hot-rolled bar, billet, plate, or bar flats, or it may be an intermediate processing step for products such as cold-finished bar, rod, and wire, and cold-rolled sheet and strip. Typical properties of hot processed products and of hot processed and annealed products are different from those of either cast or cold reduced products. Hot processed products tend to have coarser grain sizes than cold reduced products. Cold Reduced Products. When strained at ambient temperatures, all stainless steels tend to work harden, as shown in Fig. 13 . Because recrystallization does not occur during cold working, the final properties of thermally treated products depend on: • Amount of cold reduction (which helps determine the number of potential recrystallization sites) • Type of mill thermal treatment (subcritical annealing, normalizing, or solution treatment) • Time at any given temperature


Page 1349



01 Sep 2005

Wrought products that have been cold reduced and annealed generally have finer grain sizes, which produce higher strengths than hot processed products. Cold reduced products sometimes exhibit greater differences between transverse and longitudinal properties than hot processed products. Fig. 13 Typical effect of cold rolling on the tensile strength of selected stainless steels

Cold finishing is generally done to improve dimensional tolerances or surface finish or to raise mechanical strength. Cold-finished products⎯whether they have been previously hot worked and annealed or have been hot worked, cold worked, and annealed⎯have higher mechanical strength and slightly lower ductility than their process-annealed counterparts.

Physical Properties There are relatively few applications for stainless steels in which physical properties are the determining factors in selection. However, there are many applications in which physical properties are important in product design. For instance, stainless steels are used for many elevated-temperature applications, often in conjunction with steels of lesser alloy content. Because austenitic stainless steels have higher coefficients of thermal expansion and lower thermal conductivities than carbon and alloy steels, these characteristics must be taken into account in the design of stainless steel-to-carbon steel or stainless steel-to-alloy steel products such as heat exchangers. In such products, differential thermal expansion imposes stresses on the unit that would not be present were the unit made entirely of carbon or alloy steel; also, if the heat-transfer surface is made of stainless steel, it must be larger than if it were made of carbon or alloy steel. Typical physical properties of selected grades of annealed wrought stainless steels are given in Table 21 . Physical properties may vary slightly with product form and size, but such variations are usually not of critical importance to the application. Table 21 Typical physical properties of wrought stainless steels, annealed condition Mean CTE from 0 °C (32 °F) to:

Type

100 °C 315 °C 538 °C Elastic (212 (600 °F) (1000 °F) modulu °F) µm/m · µm/m · µm/m · s °C °C °C UNS Density GPa 3 (106 (µin./in. (µin./in. · (µin./in. · numb g/cm · °F) °F) °F) er (lb/in.3 psi)

Thermal conductivity

at 100 °C (212 °F) W/m · K (Btu/ft · h · °F)

Specific at 500 °C heat(a) Electric Magnet (932 °F) J/kg · al ic K resistivi permeW/m · K ty, (Btu/ft · h (Btu/lb ability( · °F) nΩ · m · °F) b)

Melting range °C (°F)

201

S2010 7.8 0 (0.28)

197 (28.6)

15.7 (8.7)

17.5 (9.7)

18.4 (10.2)

16.2 (9.4)

21.5 (12.4)

500 (0.12)

690

1.02

1400-1450 (2550−2650)

202

S2020 7.8 0 (0.28)

...

17.5 (9.7)

18.4 (10.2)

19.2 (10.7)

16.2 (9.4)

21.6 (12.5)

500 (0.12)

690

1.02

1400−1450 (2550−2650)

205

S2050 7.8 0 (0.28)

197 (28.6)

...

17.9 (9.9)

19.1 (10.6)

...

...

500 (0.12)

...

...

...

301

S3010 8.0 0 (0.29)

193 (28.0)

17.0 (9.4)

17.2 (9.6)

18.2 (10.1)

16.2 (9.4)

21.5 (12.4)

500 (0.12)

720

1.02

1400−1420 (2550−2590)

302

S3020 8.0 0 (0.29)

193 (28.0)

17.2 (9.6)

17.8 (9.9)

18.4 (10.2)

16.2 (9.4)

21.5 (12.4)

500 (0.12)

720

1.02

1400−1420 (2550−2590)

302B

S3021 8.0 5 (0.29)

193 (28.0)

16.2 (9.0)

18.0 (10.0)

19.4 (10.8)

15.9 (9.2)

21.6 (12.5)

500 (0.12)

720

1.02

1375−1400 (2550−2550)

303

S3030

193

17.2

17.8 (9.9)

18.4

16.2 (9.4)

21.5

500

720

1.02

1400−1450

8.0


Page 1350


0


(0.29)

(28.0)

(9.6)

(10.2)

304

S3040 8.0 0 (0.29)

193 (28.0)

17.2 (9.6)

17.8 (9.9)

18.4 (10.2)

304L

S3040 8.0 3 (0.29)

...

...

...

302Cu

S3043 8.0 0 (0.29)

193 (28.0)

17.2 (9.6)

304N

S3045 8.0 1 (0.29)

196 (28.5)

305

S3050 8.0 0 (0.29)

308

01 Sep 2005

(12.4)

(0.12)

16.2 (9.4)

21.5 (12.4)

500 (0.12)

720

1.02

1400−1450 (2550−2650)

...

...

...

...

...

1.02

1400−1450 (2550−2650)

17.8 (9.9)

...

11.2 (6.5)

21.5 (12.4)

500 (0.12)

720

1.02

1400−1450 (2550−2650)

...

...

...

...

...

500 (0.12)

720

1.02

1400−1450 (2550−2650)

193 (28.0)

17.2 (9.6)

17.8 (9.9)

18.4 (10.2)

16.2 (9.4)

21.5 (12.4)

500 (0.12)

720

1.02

1400−1450 (2550−2650)

S3080 8.0 0 (0.29)

193 (28.0)

17.2 (9.6)

17.8 (9.9)

18.4 (10.2)

15.2 (8.8)

21.6 (12.5)

500 (0.12)

720

...

1400−1420 (2550−2590)

309

S3090 8.0 0 (0.29)

200 (29.0)

15.0 (8.3)

16.6 (9.2) 17.2 (9.6) 15.6 (9.0)

18.7 (10.8)

500 (0.12)

780

1.02

1400−1450 (2550−2650)

310

S3100 8.0 0 (0.29)

200 (29.0)

15.9 (8.8)

16.2 (9.0) 17.0 (9.4) 14.2 (8.2)

18.7 (10.8)

500 (0.12)

780

1.02

1400−1450 (2550−2650)

314

S3140 7.8 0 (0.28)

200 (29.0)

...

15.1 (8.4)

20.9 (12.1)

500 (0.12)

770

1.02

...

316

S3160 8.0 0 (0.29)

193 (28.0)

15.9 (8.8)

21.5 (12.4)

500 (0.12)

740

1.02

1375−1400 (2500−2550)

316L

S3160 8.0 3 (0.29)

...

...

...

...

...

...

...

...

1.02

1375−1400 (2500−2550)

316N

S3165 8.0 1 (0.29)

196 (28.5)

...

...

...

...

...

500 (0.12)

740

1.02

1375−1400 (2500−2550)

317

S3170 8.0 0 (0.29)

193 (28.0)

15.9 (8.8)

21.5 (12.4)

500 (0.12)

740

1.02

1375−1400 (2500−2550)

317L

S3170 8.0 3 (0.29)

200 (29.0)

16.5 (9.2)

...

18.1 (10.1)

14.4 (8.3)

...

500 (0.12)

790

...

1375−1400 (2500−2550)

321

S3210 8.0 0 (0.29)

193 (28.0)

16.6 (9.2)

17.2 (9.6)

18.6 (10.3)

16.1 (9.3)

22.2 (12.8)

500 (0.12)

720

1.02

1400−1425 (2550−2600)

329

S3290 7.8 0 (0.28)

...

...

...

...

...

...

460 (0.11)

750

...

...

330

N083 30

8.0 (0.29)

196 (28.5)

14.4 (8.0)

16.0 (8.9) 16.7 (9.3)

...

...

460 (0.11)

1020

1.02

1400−1425 (2550−2600)

347

S3470 8.0 0 (0.29)

193 (28.0)

16.6 (9.2)

17.2 (9.6)

18.6 (10.3)

16.1 (9.3)

22.2 (12.8)

500 (0.12)

730

1.02

1400−1425 (2550−2600)

384

S3840 8.0 0 (0.29)

193 (28.0)

17.2 (9.6)

17.8 (9.9)

18.4 (10.2)

16.2 (9.4)

21.5 (12.4)

500 (0.12)

790

1.02

1400−1450 (2550−2650)

405

S4050 7.8 0 (0.28)

200 (29.0)

10.8 (6.0)

11.6 (6.4) 12.1 (6.7)

27.0 (15.6)

...

460 (0.11)

600

...

1480−1530 (2700−2790)

409

S4090 7.8 0 (0.28)

...

11.7 (6.5)

...

...

...

...

...

1480−1530 (2700−2790)

410

S4100 7.8 0 (0.28)

200 (29.0)

9.9 (5.5)

11.4 (6.3) 11.6 (6.4)

24.9 (14.4)

28.7 (16.6)

460 (0.11)

570

700−10 00

1480−1530 (2700−2790)

414

S4140 7.8 0 (0.28)

200 (29.0)

10.4 (5.8)

11.0 (6.1) 12.1 (6.7)

24.9 (14.4)

28.7 (16.6)

460 (0.11)

700

...

1425−1480 (2600−2700)

416

S4160 7.8 0 (0.28)

200 (29.0)

9.9 (5.5)

11.0 (6.1) 11.6 (6.4)

24.9 (14.4)

28.7 (16.6)

460 (0.11)

570

700−10 00

1480−1530 (2700−2790)

420

S4200 7.8 0 (0.28)

200 (29.0)

10.3 (5.7)

10.8 (6.0) 11.7 (6.5)

24.9 (14.4)

...

460 (0.11)

550

...

1450−1510 (2650−2750)

422

S4220 7.8 0 (0.28)

...

11.2 (6.2)

11.4 (6.3) 11.9 (6.6)

23.9 (13.8)

27.3 (15.8)

460 (0.11)

...

...

1470−1480 (2675−2700)

...

17.5 (10.1)

16.2 (9.0) 17.5 (9.7) 16.2 (9.4)

16.2 (9.0) 17.5 (9.7) 16.2 (9.4)

...


...

(2550−2590)

Page 1351



429

S4290 7.8 0 (0.28)

200 (29.0)

10.3 (5.7)

430

S4300 7.8 0 (0.28)

200 (29.0)

10.4 (5.8)

430F

S4302 7.8 0 (0.28)

200 (29.0)

431

S4310 7.8 0 (0.28)

434

...

25.6 (14.8)

...

460 (0.11)

590

...

1450−1510 (2650−2750)

11.0 (6.1) 11.4 (6.3)

26.1 (15.1)

26.3 (15.2)

460 (0.11)

600

600−11 00

1425−1510 (2600−2750)

10.4 (5.8)

11.0 (6.1) 11.4 (6.3)

26.1 (15.1)

26.3 (15.2)

460 (0.11)

600

...

1425−1510 (2600−2750)

200 (29.0)

10.2 (5.7)

12.1 (6.7)

20.2 (11.7)

...

460 (0.11)

720

...

...

S4340 7.8 0 (0.28)

200 (29.0)

10.4 (5.8)

11.0 (6.1) 11.4 (6.3)

...

26.3 (15.2)

460 (0.11)

600

600−11 00

1425−1510 (2600−2750)

436

S4360 7.8 0 (0.28)

200 (29.0)

9.3 (5.2)

23.9 (13.8)

26.0 (15.0)

460 (0.11)

600

600−11 00

1425−1510 (2600−2750)

439

S4303 7.7 5 (0.28)

200 (29.0)

10.4 (5.8)

24.2 (14.0)

...

460 (0.11)

630

...

...

440A

S4400 7.8 2 (0.28)

200 (29.0)

10.2 (5.7)

...

...

24.2 (14.0)

...

460 (0.11)

600

...

1370−1480 (2500−2700)

440C

S4400 7.8 4 (0.28)

200 (29.0)

10.2 (5.7)

...

...

24.2 (14.0)

...

460 (0.11)

600

...

1370−1480 (2500−2700)

444

S4440 7.8 0 (0.28)

200 (29.0)

10.0 (5.6)

10.6 (5.9) 11.4 (6.3)

26.8 (15.5)

...

420 (0.10)

620

...

...

446

S4460 7.5 0 (0.27)

200 (29.0)

10.4 (5.8)

10.8 (6.0) 11.2 (6.2)

20.9 (12.1)

24.4 (14.1)

500 (0.12)

670

400−70 0

1425−1510 (2600−2750)

PH 13-8 Mo

S1380 7.8 0 (0.28)

203 (29.4)

10.6 (5.9)

11.2 (6.2) 11.9 (6.6) 14.0 (8.1)

22.0 (12.7)

460 (0.11)

1020

...

1400−1440 (2560−2625)

15-5 PH

S1550 7.8 0 (0.28)

196 (28.5)

10.8 (6.0)

11.4 (6.3)

...

17.8 (10.3)

23.0 (13.1)

420 (0.10)

770

95

1400−1440 (2560−2625)

17-4 PH

S1740 7.8 0 (0.28)

196 (28.5)

10.8 (6.0)

11.6 (6.4)

...

18.3 (10.6)

23.0 (13.1)

460 (0.11)

800

95

1400−1440 (2560−2625)

17-7 PH

S1770 7.8 0 (0.28)

204 (29.5)

11.0 (6.1)

11.6 (6.4)

...

16.4 (9.5)

21.8 (12.6)

460 (0.11)

830

...

1400−1440 (2560−2625)

...

...

01 Sep 2005

...

...

11.0 (6.1) 11.4 (6.3)

CTE, coefficient of thermal expansion.(a) At 0 to 100 °C (32 to 212 °F). (b) Approximate values

Corrosion Properties Stainless steels are susceptible to several forms of localized corrosive attack. The avoidance of such localized corrosion is the focus of much of the effort involved in selecting stainless steel. Furthermore, the corrosion performance of stainless steels can be strongly affected by practices of design, fabrication, surface conditioning, and maintenance. The selection of a grade of stainless steel for a particular application involves the consideration of many factors, but always begins with corrosion resistance. It is first necessary to characterize the probable service environment. It is not enough to consider only the design conditions. It is also necessary to consider the reasonably anticipated excursions or upsets in service conditions. The suitability of various grades can be estimated from laboratory tests or from documentation of field experience in comparable environments. Once the grades with adequate corrosion resistance have been identified, it is then appropriate to consider mechanical properties, ease of fabrication, the types and degree of risk present in the application, the availability of the necessary product forms, and cost. Mechanism of Corrosion Resistance The mechanism of corrosion protection for stainless steels differs from that for carbon steels, alloy steels, and most other metals. In these other cases, the formation of a barrier of true oxide separates the metal from the surrounding atmosphere. The degree of protection afforded by such an oxide is a function of the thickness of the oxide layer, its continuity, its coherence and adhesion to the metal, and the diffusivities of oxygen and metal in the oxide. In high-temperature oxidation, stainless steels use a generally similar model for corrosion protection. However, at low temperatures, stainless steels do not form a layer of true oxide. Instead, a passive film is formed. One mechanism that has been suggested is the formation of a film of hydrated oxide, but there is not total agreement on the nature of the oxide complex on the metal surface. However, the oxide film should be continuous, nonporous, insoluble, and self-healing if broken in the presence of oxygen. Passivity exists under certain conditions for particular environments. The range of conditions over which passivity can be maintained depends on the precise environment and on the family and composition ofthe stainless steel. When conditions are


Page 1352



01 Sep 2005

favorable for maintaining passivity, stainless steels exhibit extremely low corrosion rates. If passivity is destroyed under conditions that do not permit the restoration of the passive film, stainless steel will corrode much like a carbon or low-alloy steel. The presence of oxygen is essential to the corrosion resistance of a stainless steel. The corrosion resistance of stainless steel is at its maximum when the steel is boldly exposed and the surface is maintained free of deposits by a flowing bulk environment. Covering a portion of the surface, for example, by biofouling, painting, or installing a gasket, produces an oxygen-depleted region under the covered region. The oxygen-depleted region is anodic relative to the well-aerated boldly exposed surface, and a higher level of alloy content in the stainless steel is required to prevent corrosion. With appropriate grade selection, stainless steel will perform for very long times with minimal corrosion, but an inadequate grade can corrode and perforate more rapidly than a plain carbon steel will fail by uniform corrosion. The selection of the appropriate grade of stainless steel, then, is a balancing of the desire to minimize cost and the risk of corrosion damage by excursions of environmental conditions during operation or downtime. Confusion exists regarding the meaning of the term passivation. It is not necessary to chemically treat a stainless steel to obtain the passive film; the film forms spontaneously in the presence of oxygen. Most frequently, the function of passivation is to remove free iron, oxides, and other surface contamination. For example, in the steel mill, the stainless steel may be pickled in an acid solution, often a mixture of nitric and hydrofluoric acids (HNO3-HF), to remove oxides formed in heat treatment. Once the surface is cleaned and the bulk composition of the stainless steel is exposed to air, the passive film forms immediately. Effects of Composition Chromium is the one element essential in forming the passive film. Other elements can influence the effectiveness of chromium in forming or maintaining the film, but no other element can, by itself, create the properties of stainless steel. Chromium. The film is first observed at about 10.5% Cr, but it is rather weak at this composition and affords only mild atmospheric protection. Increasing the chromium content to 17 to 20%, typical of the austenitic stainless steels, or to 26 to 29%, as possible in the newer ferritic stainless steels, greatly increases the stability of the passive film. However, higher chromium may adversely affect mechanical properties, fabricability, weldability, or suitability for applications involving certain thermal exposures. Therefore, it is often more efficient to improve corrosion resistance by altering other elements, with or without some increase in chromium. Nickel, in sufficient quantities, will stabilize the austenitic structure; this greatly enhances mechanical properties and fabrication characteristics. Nickel is effective in promoting repassivation, especially in reducing environments. Also, it is particularly useful in resisting corrosion in mineral acids. Increasing nickel content to about 8 to 10% decreases resistance to stress-corrosion cracking (SCC), but further increases begin to restore SCC resistance. Resistance to SCC in most service environments is achieved at about 30% Ni. In the newer ferritic grades, in which the nickel addition is less than that required to destabilize the ferritic phase, there are still substantial effects. In this range, nickel increases yield strength, toughness, and resistance to reducing acids, but makes the ferritic grades susceptible to SCC in concentrated magnesium chloride (MgCl2)solutions. Most chloride cracking testing has been carried out in accelerated test media such as boiling MgCl2 solution (boiling point: 154 °C, or 309 °F) (Ref 32, 33, 34). All austenitic stainless steels are susceptible to chloride cracking (Fig. 14 ). It is noteworthy, however, that the higher-nickel types 310 and 314 were appreciably more resistant than the others. Although this solution causes rapid cracking, it does not necessarily simulate the cracking observed in field applications. Fig. 14 Relative SCC behavior of austenitic stainless steels in boiling magnesium chloride. Source: Ref 35


Page 1353



01 Sep 2005

Manganese in moderate quantities and in association with nickel additions will perform many of the functions attributed to nickel. However, total replacement of nickel by manganese is not practical. Very high-manganese steels have some unusual and useful mechanical properties, such as resistance to galling. Manganese interacts with sulfur in stainless steels to form manganese sulfides. The morphology and composition of these sulfides can have substantial effects on corrosion resistance, especially pitting resistance. Molybdenum in combination with chromium is very effective in terms of stabilizing the passive film in the presence of chlorides. Molybdenum is especially effective in increasing resistance to the initiation of pitting and crevice corrosion. Carbon is useful to the extent that it permits hardenability by heat treatment, which is the basis of the martensitic grades, and provides strength in the high-temperature applications of stainless steels. In all other applications, carbon is detrimental to corrosion resistance through its reaction with chromium. In the ferritic grades, carbon is also extremely detrimental to toughness. Nitrogen is beneficial to austenitic stainless steels in that it enhances pitting resistance, retards the formation of the chromium-molybdenum σ phase, and strengthens the steel. Nitrogen is essential in the newer duplex grades for increasing the austenite content, diminishing chromium and molybdenum segregation, and raising the corrosion resistance of the austenitic phase. Nitrogen is highly detrimental to the mechanical properties of the ferritic grades and must be treated as comparable to carbon when a stabilizing element is added to the steel. Forms of Corrosion of Stainless Steels The various forms of corrosive attack will be briefly discussed in this section. Detailed information on each of these forms of corrosion is available in the Section "Forms of Corrosion" in Corrosion, Volume 13 of ASM Handbook, formerly 9th Edition Metals Handbook. General (uniform) corrosion of a stainless steel suggests an environment capable of stripping the passive film from the surface and preventing repassivation. Such an occurrence could indicate an error in grade selection. An example is the exposure of a lower-chromium ferritic stainless steel to moderate concentration of hot sulfuric acid (H2SO4). Galvanic corrosion results when two dissimilar metals are in electrical contact in a corrosive medium. As a highly corrosion-resistant metal, stainless steel can act as a cathode when in contact with a less noble metal, such as steel. The corrosion of steel parts, for example, steel bolts in a stainless steel construction, can be a significant problem. However, the effect can be used in a beneficial way for protecting critical stainless steel components within a larger steel construction. In the case of stainless steel connected to a more noble metal, the active-passive condition of the stainless steel must be considered. If the stainless steel is passive in the environment, galvanic interaction with a more noble metal is unlikely to produce significant corrosion. If the stainless steel is active or only marginally passive, galvanic interaction with a more noble metal will probably produce sustained rapid corrosion of the stainless steel without repassivation. The most important aspect of galvanic interaction for stainless steels is the need to select fasteners and weldments of adequate corrosion resistance relative to the bulk material, which is likely to have a much larger exposed area. Pitting is a localized attack that can produce the penetration of a stainless steel with almost negligible weight loss to the total structure. Pitting is associated with a local discontinuity of the passive film. It can be a mechanical imperfection, such as an inclusion or surface damage, or it can be a local chemical breakdown of the film. Chloride is the most common agent for the


Page 1354



01 Sep 2005

initiation of pitting. Once a pit is formed, it in effect becomes a crevice; the local chemical environment is substantially more aggressive than the bulk environment. Therefore, very high flow rates over a stainless steel surface tend to reduce pitting corrosion; a high flow rate prevents the concentration of corrosive species in the pit. The stability of the passive film with respect to resistance to pitting initiation is controlled primarily by chromium and molybdenum. Minor alloying elements can also have an important effect by influencing the amount and type of inclusions (for example, sulfides) in the steel that can act as pitting sites. Pitting initiation can also be influenced by surface condition, including the presence of deposits, and by temperature. For a particular environment, a grade of stainless steel may be characterized by a single temperature, or a very narrow range of temperatures, above which pitting will initiate and below which pitting will not initiate. It is therefore possible to select a grade that will not be subject to pitting attack if the chemical environment and temperature do not exceed the critical levels. If the range of operating conditions can be accurately characterized, a meaningful laboratory evaluation is possible. The formation of deposits in service can reduce the pitting temperature. Although chloride is known to be the primary agent of pitting attack, it is not possible to establish a single critical chloride limit for each grade. The corrosivity of a particular concentration of chloride solution can be profoundly affected by the presence or absence of various other chemical species that may accelerate or inhibit corrosion. Chloride concentration may increase where evaporation or deposits occur. Because of the nature of pitting attack⎯rapid penetration with little total weight loss⎯it is rare for any significant amount of pitting to be acceptable in practical application. Crevice corrosion can be considered a severe form of pitting. Any crevice, whether the result of a metal-to-metal joint, a gasket, fouling, or deposits, tends to restrict oxygen access, resulting in attack. In practice, it is extremely difficult to prevent all crevices, but every effort should be made to do so. Higher-chromium, and especially higher-molybdenum, grades are more resistant to crevice attack. Just as there is a critical pitting temperature for a particular environment, there is a critical crevice temperature. This temperature is specific to the geometry and nature of the crevice and to the precise corrosion environment for each grade. The critical crevice temperature can be useful in selecting an adequately resistant grade for a particular application. Intergranular corrosion is a preferential attack at the grain boundaries of a stainless steel. It is generally the result of sensitization. This condition occurs when a thermal cycle leads to grain-boundary precipitation of a carbide, nitride, or intermetallic phase without providing sufficient time for chromium diffusion to fill the locally depleted region. A grain-boundary precipitate is not the point of attack; instead, the low-chromium region adjacent to the precipitate is susceptible. Sensitization is not necessarily detrimental unless the grade is to be used in an environment capable of attacking the region. For example, elevated-temperature applications for stainless steel can operate with sensitized steel, but concern for intergranular attack must be given to possible corrosion during downtime when condensation might provide a corrosive medium. Because chromium provides corrosion resistance, sensitization also increases the susceptibility of chromium-depleted regions to other forms of corrosion, such as pitting, crevice corrosion, and stress-corrosion cracking (SCC). The thermal exposures required to sensitize a steel can be relatively brief, as in high-temperature service. Stress-corrosion cracking is a corrosion mechanism in which the combination of a susceptible alloy, sustained tensile stress, and a particular environment leads to cracking of the metal. Stainless steels are particularly susceptible to SCC in chloride environments; temperature and the presence of oxygen tend to aggravate chloride SCC of stainless steels. Most ferritic and duplex stainless steels are either immune or highly resistant to SCC. All austenitic grades, especially AISI types 304 and 316, are susceptible to some degree. The highly alloyed austenitic grades are resistant to sodium chloride (NaCl) solutions, but crack readily in MgCl2 solutions. Although some localized pitting or crevice corrosion probably precedes SCC, the amount of pitting or crevice attack may be so small that it is undetectable. Stress corrosion is difficult to detect while in progress, even when pervasive, and can lead to rapid catastrophic failures of pressurized equipment. It is difficult to alleviate the environmental conditions that lead to SCC. The level of chlorides required to produced stress-corrosion cracking is very low. In operation, there can be evaporative concentration or a concentration in the surface film on a heat-rejecting surface. Temperature is often a process parameter, as in the case of a heat exchanger. Tensile stress is one parameter that might be controlled. However, the residual stresses associated with fabrication, welding, or thermal cycling, rather than design stresses, are often responsible for SCC, and even stress-relieving heat treatments do not completely eliminate these residual stresses. Erosion-Corrosion. Corrosion of a metal or alloy can be accelerated when there is an abrasive removal of the protective oxide layer. This form of attack is especially significant when the thickness of the oxide layer is an important factor in determining corrosion resistance. In the case of a stainless steel, erosion of the passive film can lead to some acceleration of attack. Oxidation. Because of their high chromium contents, stainless steels tend to be very resistant to oxidation. Important factors to be considered in the selection of stainless steel grades for high-temperature service are the stability of the composition and microstructure upon thermal exposure and the adherence of the oxide scale upon thermal cycling. Because many of the stainless steels used for high temperatures are austenitic grades with relatively high nickel contents, it is also necessary to be alert to the possibility of sulfidation attack. Corrosion in Specific Environments The selection of a suitable stainless steel for a specific environment requires consideration of several criteria. The first is corrosion resistance. Alloys are available that provide resistance to mild atmospheres (for example, type 430) or to many food-processing environments (for example, type 304 stainless). Chemicals and more severe corrodents require type 316 or a


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more highly alloyed material, such as 20Cb-3 (UNS N08020). Factors that affect the corrosivity of an environment include the concentration of chemical species, pH, aeration, flow rate (velocity), impurities (such as chlorides), and temperature, including effects from heat transfer. The second criterion is mechanical properties, or strength. High-strength materials often sacrifice resistance to some form of corrosion, particularly SCC. Third, fabrication must be considered, including such factors as the ability of the steel to be machined, welded, or formed. Resistance of the fabricated article to the environment must be considered, for example, the ability of the material to resist attack in crevices that cannot be avoided in the design. Fourth, total cost must be estimated, including initial alloy price, installed cost, and the effective life expectancy of the finished product. Finally, consideration must be given to product availability. Many applications for stainless steels, particularly those involving heat exchangers, can be analyzed in terms of a process side and a water side. The process side is usually a specific chemical combination that has its own requirements for a stainless steel grade. The water side is common in many applications. This section will discuss the corrosivity of various environments for stainless steels. Atmospheric Corrosion. The atmospheric contaminants most often responsible for the rusting of structural stainless steels are chlorides and metallic iron dust. Chloride contamination may originate from the calcium chloride (CaCl2 used to make concrete or from exposure in marine or industrial locations. Iron contamination may occur during fabrication or erection of the structure. Contamination should be minimized, if possible. The corrosivity of different atmospheric exposures can vary greatly and can dictate application of different grades of stainless steel. Rural atmospheres, uncontaminated by industrial fumes or coastal salt, are extremely mild in terms of corrosivity for stainless steel, even in areas of high humidity. Industrial or marine environments can be considerably more severe. Table 22 demonstrates that resistance to staining can depend on the specific exposure. For example, several 300-series stainless steels showed no rust during long-term exposures in New York City. On the other hand, staining was observed after much shorter exposures at Niagara Falls in a severe industrial-chemical environment near plants producing chlorine and HCl. Table 22 Atmospheric corrosion of austenitic stainless steels at two industrial sites New York City (industrial)

Niagara Falls (industrial-chemical)

Exposure time, years

Specimen surface evaluation

Exposure time, years

302

5

Free from rust stains

~25%) sections should be fully annealed and quenched whenever applications to the alloy present the possibility of stress-corrosion cracking in the service environment. Full annealing is


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conducted in the temperature range of 1010 to 1100 °C (1850 to 2010 °F), followed by a rapid cooling. Hot forming operations are usually performed in the temperature range of 980 to 1260 °C (1800 to 2300 °F); the preferred temperature range is dictated by the specific alloy composition. The temperature range of about 370 to 925 °C (700 to 1700 °F) should be avoided to preclude the precipitation of such deleterious phases as σ and α′. These phases can adversely affect mechanical properties and corrosion resistance. Duplex stainless steels, especially in light gages, can be made to exhibit superplastic behavior at elevated temperatures. This property may be used to advantage in certain operations involving the continuous line annealing of strip and sheet. Comparison With Carbon Steel. The curves for 1008 low-carbon steel are included in Fig. 24 as a reference for the evaluation of stainless steels. The decrease in formability of 1008 steel with cold work appears to fall between that of types 409/430 and that of the more formable type 301. Figure 24 also shows that cold work does not increase the strength of 1008 as rapidly as it does that of type 301 and the ferritic alloys. Stress-Strain Relationships. Figure 25 shows load-elongation curves for six types of stainless steel: four austenitic (202, 301, 302, and 304), one martensitic (410), and one ferritic (430). The figure also shows that the type of failure in the cup drawing of the austenitic types was different from that of types 410 and 430, as shown in Fig. 25 . The austenitic types broke in a fairly clean line near the punch nose radius, almost as though the bottom of the drawn cup were blanked out; types 410 and 430 broke in the sidewall in sharp jagged lines, showing extreme brittleness as a result of the severe cold work. Fig. 25 Comparison of ductility of six stainless steels and of the types of failure resulting from deep drawing

As suggested by the data in Fig. 25 , the power required to form type 301 exceeds that required by the other austenitic alloys. In addition, type 301 will develop maximum elongation before failing. Types 410 and 430 require considerably less power to form, but fail at comparatively low elongation levels. Power requirements for forming stainless steel, because of the high yield strength, are greater than those for low-carbon steel; generally, twice as much power is used in forming stainless steel. Because the austenitic steels work harden rapidly in cold-forming operations, the need for added power after the start of initial deformation is greater than that for the ferritic steels.


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The ferritic steels behave much like plain carbon steels once deformation begins, although higher power is also needed to start plastic deformation. Forgeability Stainless steels, based on forging pressure and load requirements, are considerably more difficult to forge than carbon or low-alloy steels, primarily because of the greater strength of stainless steels at elevated temperatures and the limitations on the maximum temperatures at which stainless steels can be forge without incurring microstructural damage. Forging load requirements and forgeability vary widely among stainless steels of different types and compositions; the most difficult alloys to forge are those with the greatest strength at elevated temperatures. Ingot Breakdown. In discussing the forgeability of the stainless steels, it is critical to understand the types of primary mill practices available to the user of semifinished billet or bloom product. Primary Forging and Ingot Breakdown. Most stainless steel ingots destined for the forge shop are melted by the electric furnace argon-oxygen decarburization process. They usually weigh between 900 and 13,500 kg (2000 to 30,000 lb), depending on the shop and the size of the finished piece. Common ingot shapes are round, octagonal, and fluted; less common ingot shapes include squares. Until recently, all of these ingots would have been top poured. Increasing numbers of producers are switching to the bottom-poured ingot process. This process is slightly more expensive to implement in the melt shop, but it more than pays for itself in extended mold life and greatly improved ingot surface. Some stainless steel grades used in the aircraft and aerospace industries are double melted. The first melt is done with the electric furnace and argon-oxygen decarburization, and these electrodes are then remelted by a vacuum arc remelting (VAR) or electroslag remelting (ESR) process. This remelting under a vacuum (VAR) or a slag (ESR) tends to give a much cleaner product with better hot workability. For severe forging applications, the use of remelt steels can sometimes be a critical factor in producing acceptable parts. These double-melted ingots are round in shape and vary in diameter from 460 to 915 mm (18 to 36 in.), and in some cases, they weigh in excess of 11,000 kg (25,000 lb). The breakdown of ingots is usually done on large hydraulic presses (13,500 kN, or 1500 tonf). A few shops, however, still use large hammers, and the four-hammer radial forging machine is being used increasingly for ingot breakdown. Heating is the single most critical step in the initial forging of ingots. The size of the ingot and the grade of the end-product stainless steel dictate the practice necessary to reduce thermal shock and to avoid unacceptable segregation levels. It is essential to have accurate and programmable control of the furnaces used to heat stainless steel ingots and large blooms. Primary forging or breakdown of an ingot is usually achieved using flat dies. However, some forgers work the ingot down as a round, using V or swage dies. Because of the high hot hardness of stainless steel and the narrow range of working temperatures for these alloys, light reductions, or saddening is the preferred initial step in the forging of the entire surface of the ingot. Saddening is an operation in which an ingot is given a succession of light reductions in a press or rolling mill or under a hammer in order to break down the skin and overcome the initial fragility due to a coarse crystalline structure preparatory to reheating prior to heavier reductions. After the initial saddening of the ingot surface is complete, normal reductions of 50 to 100 mm (2 to 4 in.) can be taken. If the chemistry of the heat is in accordance with specifications and if heating practices have been followed and minimum forging temperatures observed, no problems should be encountered in making the bloom and other semifinished product. If surface tears occur, the forging should be stopped, and the workpiece conditioned. Some forgers use hot powder scarfing, but this presents environmental problems. The most common method is to grind out the defect. The ferritic, austenitic, and nitrogen-strengthened austenitic stainless steels can be air cooled, ground, and reheated for reforging. the martensitic and precipitation-hardening grades must be slow cooled and overaged before grinding and reheating. The ingot surface is important, and many producers find it advantageous to grind the ingots before forging to ensure good starting surfaces. Billet and Bloom Product. Forgers buy bars, billets, or blooms of stainless steel for subsequent forging on hammers and presses. Forged stainless steel billet and bloom products tend to have better internal integrity than rolled product, especially with larger-diameter sections (>180 mm, or 7 in.). Correctly conditioned billet and bloom product should yield acceptable finished forgings if good heating practices are followed and if attention is paid to the minimum temperature requirements. Special consideration must be given to sharp corners and thin sections, because these tend to cool off very rapidly. Precautions should be taken when forging precipitation-hardening or nitrogen-strengthened austenitic grades. Closed-Die Forgeability. The relative forging characteristics of stainless steels can be most easily depicted through examples of closed-die forgings. The forgeability trends these examples establish can be interpreted in light of the grade, type of part, and forging method to be used. Stainless steels of the 300 and 400 series can be forged into any of the hypothetical parts illustrated in Fig. 26 . However, the forging of stainless steel into shapes equivalent in severity to part 3 may be prohibited by shortened die life (20 to 35% of that obtained in forging such a shape from carbon or low-alloy steel) and by the resulting high cost. For a given shape, die life is shorter in forging stainless steel than in forging carbon or low-alloy steel. Fig. 26 Three degrees of forging severity. Dimensions given in inches


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Forgings of mild severity, such as part 1 in Fig. 26 , can be produced economically from any stainless steel with a single heating and about five blows. Forgings approximating the severity of part 2 can be produced from any stainless steel with a single heating and about ten blows. For any type of stainless steel, die life in the forging of part 1 will be about twice that in the forging of part 2. Part 3 represents the maximum severity for forging all stainless steels and especially those with high strength at elevated temperature, namely, types 309, 310, 314, 316, 317, 321, and 347. Straight-chromium types 403, 405, 410, 416, 420, 430, 431, and 440 are the easiest to forge into a severe shape such as part 3 (although type 440, because of its high carbon content, would be the least practical). Types 201, 301, 302, 303, and 304 are intermediate between the two previous groups. One forge shop has reported that part 3 would be practical and economical to produce in the higher-strength alloys if the center web were increased from 3.2 to 6.4 mm (1=8to 1=4in.) and if all fillets and radii were increased in size. It could then be forged with 15 to 20 blows and one reheating, dividing the number of blows about equally between the first heat and the reheat. Hot Upsetting. Stainless steel forgings of the severity represented by hypothetical parts 4, 5, and 6 in Fig. 27 can be hot upset in one blow in a steel die. However, the conditions are similar to those encountered in hot-die forging. First, with a stainless steel, die wear in the upsetting of part 6 will be several times as great as in the upsetting of part 4. Second, die wear for the forming of any shape will increase as the elevated-temperature strength of the alloy increases. Therefore, type 410, with about the lowest strength at high temperature, would be the most economical stainless steel to be formed of any of the parts, particularly part 6. Conversely, type 310 would be the least economical. Fig. 27 Three degrees of upsetting severity

Upset Reduction Versus Forging Pressure. The effect of percentage of upset reduction (upset height versus original height) on forging pressure for low-carbon steel and for type 304 stainless steel at various temperatures is shown in Fig. 28 . Temperature has a marked effect on the pressure required for any given percentage of upset, and at any given forging temperature and percentage of upset, type 304 stainless requires at least twice the pressure required for 1020 steel. Fig. 28 Effect of upset reduction on forging pressure for various temperatures. Source: Ref 61

The effects of temperature on forging pressure are further emphasized in Fig.(a) 29 . These data, based on an upset reduction of 10%, show that at 760 °C (1400 °F) type 304 stainless steel requires only half as much pressure as A-286 (an iron-base heat-resistant alloy), although the curves for forging pressure for the two metals converge at 1100 °C (2000 °F). However, at a forging temperature of 1100 °C (2000 °F), the pressure required for a 10% upset reduction on type 304 is more than twice that required for a carbon steel (1020) and about 60% more than that required for 4340 alloy steel. Differences in forgeability, based on percentage of upset reduction and forging pressure for type 304 stainless steel, 1020, and 4340 at the same temperature (980 °C, or 1800 °F), are plotted in Fig.(b) 29 . Fig. 29 Forging pressure required for upsetting versus (a) forging temperature and (b) percentage of upset reduction. Source: Ref 62


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Austenitic stainless steels are more difficult to forge than the straight-chromium types, but are less susceptible to surface defects. Most of the austenitic stainless steels can be forged over a wide range of temperatures above 925 °C (1700 °F) and because they do not undergo major phase transformation at elevated temperature, they can be forged at higher temperatures than the martensitic types (Table 33 ). Exceptions to the above statements occur when the composition of the austenitic stainless steel promotes the formation of δ-ferrite, as in the case of the 309S, 310S, or 314 grades. At temperatures above 1100 °C (2000 °F), these steels, depending on their composition, may form appreciable amounts of δ-ferrite. Figure 30 depicts these compositional effects in terms of nickel equivalent (austenitic-forming elements) and chromium equivalent. Delta-ferrite formation adversely affects forgeability, and compensation for the amount of ferrite present can be accomplished with forging temperature restrictions. Table 33 Typical forging temperature ranges of stainless steels Temperature °C

°F

UNS N08020

980−1230

1800−2245

Pyromet 355

925−1150

1700−2100

Type 440C

925−1150

1700−2100

19-9DL/19DX

870−1150

1600−2100

Types 347 and 348

925−1230

1700−2245

Type 321

925−1260

1700−2300

Type 440B

925−1175

1700−2145

Type 440A

925−1200

1700−2200

Type 310

980−1175

1800−2145

Type 310S

980−1175

1800−2145

17-4 PH

1095−1175

2000−2145

15-5 PH

1095−1175

2000−2145

13-8 Mo

1095−1175

2000−2145

Type 317

925−1260

1700−2300

Stainless steel More difficult to hot work


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Type 316L

925−1260

1700−2300

Type 316

925−1260

1700−2300

Type 309S

980−1175

1800−2145

Type 309

980−1175

1800−2145

Type 303

925−1260

1700−2300

Type 303Se

925−1260

1700−2300

Type 305

925−1260

1700−2300

Type 329

925−1095

1700−2000

Types 302 and 304

925−1260

1700−2300

Nitronic 60

1095−1175

2000−2145

Carpenter No. 10 (Type 384)

925−1230

1700−2245

Lapelloy

1040−1150

1900−2100

AMS 5616 (Greek Ascoloy)

955−1175

1750−2145

Type 431

900−1200

1650−2200

Type 414

900−1200

1650−2200

Type 420F

900−1200

1650−2200

Type 420

900−1200

1650−2200

Type 416

925−1230

1700−2245

Type 410

900−1200

1650−2200

Type 404

900−1150

1650−2100

Type 446

900−1120

1650−2050

Type 443

900−1120

1650−2050

Type 430F

815−1150

1560−2100

815−1120

1500−2050

Easier to hot work

Type 430 Source: Ref 63

Fig. 30 Revised Schaeffler (constitution) diagram used to predict the amount of δ-ferrite that will be obtained during elevated-temperature forging or welding of austenite/ferritic stainless steels. A, austenite; M, martensite. WRC, Welding Research Council. Source: Ref 64

Equally important restrictions in forging the austenitic stainless steels apply to the finishing temperatures. All but the stabilized types (321, 347, 348) and the extralow-carbon types should be finished at temperatures above the sensitizing range (~815 to 480


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°C, or 1500 to 900 °F) and cooled rapidly from 870 °C (1600 °F) to a black heat. The highly alloyed grades, such as 309, 310, and 314, are also limited with regard to finishing temperature, because of their susceptibility at lower temperatures to hot tearing and σ formation. A final annealing by cooling rapidly from about 1065 °C (1950 °F) is generally advised for nonstabilized austenitic stainless steel forgings in order to retain the chromium carbides in solid solution. Finishing temperatures for austenitic stainless steels become more critical when section sizes increase and ultrasonic testing requirements are specified. During ultrasonic examination, coarse-grain austenitic stainless steels frequently display sweep noise that can be excessive due to a coarse-grain microstructure. The degree of sound attenuation normally increases with section size and may become too great to permit detection of discontinuities. Careful control of forging conditions, including final forge reductions of at least 5%, can assist in the improvement of ultrasonic penetrability. The stabilized or extralow-carbon austenitic stainless steels, which are resistant to sensitization, are sometimes strain hardened by small reductions at temperatures well below the forging temperature. Strain hardening is usually accomplished at 535 to 650 °C (1000 to 1200 °F) (referred to as warm working or hot-cold working). When minimum hardness is required, the forgings are solution annealed. Sulfur or selenium can be added to austenitic stainless steel to improve machinability. Selenium is preferred because harmful stringers are less likely to exist. Type 321, stabilized with titanium, may also contain stringers of segregate that will open as the surface ruptures when the steel is forged. Type 347, stabilized with niobium, is less susceptible to stringer segregation and is the stabilized grade that is usually specified for forgings. When heating the austenitic stainless steels, it is especially desirable that a slightly oxidizing furnace atmosphere be maintained. A carburizing atmosphere or an excessively oxidizing atmosphere will impair corrosion resistance, either by harmful carbon pickup or by chromium depletion. In types 309 and 310, chromium depletion can be especially severe. Nitrogen-strengthened austenitic stainless steels are iron-base alloys containing chromium and manganese. Varying amounts of nickel, molybdenum, niobium, vanadium, and/or silicon are also added to achieve specific properties. Nitrogen-strengthened austenitic stainless steels provide high strength, excellent cryogenic properties and corrosion resistance, low magnetic permeability (even after cold work or subzero temperature), and higher elevated-temperature strengths, compared to the 300-series stainless steels. A partial list of these alloys includes: • UNS S24100 (Nitronic 32) ASTM XM-28. High work hardening while remaining nonmagnetic plus twice the yield strength of type 304 with equivalent corrosion resistance • UNS S24000 (Nitronic 33) ASTM XM-29. Twice the yield strength of type 304, low magnetic permeability after severe cold work, high resistance to wear and galling compared to standard austenitic stainless steels, and good cryogenic properties • UNS S21904 (Nitronic 40) ASTM XM-11. Twice the yield strength of type 304 with good corrosion resistance, low magnetic permeability after severe cold working, and good cryogenic properties • UNS S20910 (Nitronic 50) ASTM XM-19. Corrosion resistance greater than that of type 316L with twice the yield strength, good elevated and cryogenic properties, and low magnetic permeability after severe cold work • UNS S21800 (Nitronic 60). Galling resistance with corrosion resistance equal to that of type 304 and twice the yield strength and good oxidation resistance A forgeability comparison, as defined by dynamic hot hardness, is provided in Fig. 31 . Fig. 31 Comparative dynamic hot hardness versus temperature (forgeability) for various ferrous alloys


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Martensitic stainless steels have high hardenability to the extent that they are generally air hardened. Therefore, precautions must be taken when cooling forgings of martensitic steels, especially those with high carbon content, in order to prevent cracking. The martensitic alloys are generally cooled slowly to about 595 °C (1100 °F), either by burying in an insulating medium or by temperature equalizing in a furnace. Direct water sprays, such as might be employed to cool dies, should be avoided, because they would cause cracking of the forging. Forgings of the martensitic steels are often tempered in order to soften them for machining. They are later quench hardened and tempered. Maximum forging temperatures for these steels are low enough to avoid the formation of δ-ferrite. If δ-ferrite stringers are present at forging temperatures, cracking is likely to occur. Delta-ferrite usually forms at temperatures from 1095 to 1260 °C (2000 to 2300 °F). Care must be exercised to keep the temperature below this level during forging and to avoid rapid metal movement that might result in local overheating. Surface decarburization, which promotes ferrite formation, must be minimized. The δ-ferrite formation temperature decreases with increasing chromium content, and small amounts of δ-ferrite reduce forgeability significantly. As the δ-ferrite increases above about 15% (Fig. 30 ), forgeability improves gradually until the structure becomes entirely ferritic. Finishing temperatures are limited by the allotropic transformation, which begins near 815 °C (1500 °F). However, forging of these steels is usually stopped at about 925 °C (1700 °F) because the metal is difficult to deform at lower temperatures. Sulfur or selenium can be added to type 410 to improve machinability. However, these elements can cause forging problems, particularly when they form surface stringers that open and form cracks. This can sometimes be overcome by adjusting the forging temperature or the procedure. With sulfur additions, it may be impossible to eliminate all cracking of this type. Therefore,


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selenium additions are preferred. Ferritic Stainless Steels. The ferritic straight-chromium stainless steels exhibit virtually no increase in hardness upon quenching. They will work harden during forging; the degree of work hardening depends on the temperature and the amount of metal flow. Cooling from the forging temperature is not critical. The ferritic stainless steels have a broad range of forgeability, which is restricted somewhat at higher temperature because of grain growth and structural weakness and is closely restricted in finishing temperature only for type 405. Type 405 requires special consideration because of the grain-boundary weakness resulting from the development of a small amount of austenite. The other ferritic stainless steels are commonly finished at any temperature down to 705 °C (1300 °F). For type 446, the final 10% reduction should be made below 870 °C (1600 °F) to achieve grain refinement and room-temperature ductility. Annealing after forging is recommended for ferritic steels. Precipitation-Hardening Stainless Steels. The semiaustenitic and martensitic precipitation-hardening stainless steels can be heat treated to high hardness through a combination of martensite transformation and precipitation. They are the most difficult to forge and will crack if temperature schedules are not accurately maintained. The forging range is narrow, and the steel must be reheated if the temperature falls below 980 °C (1800 °F). They have the least plasticity (greatest stiffness) at forging temperature of any of the classes and are subject to grain growth and δ-ferrite formation. Heavier equipment and a greater number of blows are required to achieve metal flow equivalent to that of the other types. During trimming, the forgings must be kept hot enough to prevent the formation of flash-like cracks. To avoid these cracks, it is often necessary to reheat the forgings slightly between the finish forging and the trimming operations. Cooling, especially the cooling of the martensitic grades, must be controlled to avoid cracking. Duplex Stainless Steels. Because of their higher strength, duplex stainless steels are generally stiff when hot worked, relative to many other stainless grades. Rolling and forging equipment must have sufficient power to reduce the material while it is in the temperature range for optimum hot workability. The surface of finished forgings can be optimized by controlling the initial surface, material composition, phase balance, and hot-working temperature. Wrought structures generally produce a better surface after reforging than does cast material after the initial forging operation. Machinability Because of the wide variety of stainless steels available, a simple characterization of their machinability can be somewhat misleading. As shown in later sections of this article, the machinability of stainless steels varies from low to very high, depending on the final choice of alloy. In general, however, stainless steels are considered more difficult to machine than other metals, such as aluminum or low-carbon steels. Stainless steels have been characterized as gummy during cutting, showing a tendency to produce long, stringy chips, which seize or form a built-up edge on the tool. This may result in reduced tool life and degraded surface finish. These general characteristics are due to the following properties possessed by stainless steels to varying degrees (Ref 65, 66) (see also the article "Machining of Stainless Steels" in Machining, Volume 16 of ASM Handbook, formerly 9th Edition Metals Handbook): • • • • •

High tensile strength (Fig. 24 ) Large spread between yield strength (YS) and ultimate tensile strength (UTS) (Fig. 24 ) High ductility and toughness High work-hardening rate (Fig. 24 ) Low thermal conductivity (Fig. 32 )

Fig. 32 Comparison of thermal conductivities for carbon steel, S30200 austenitic stainless steel, and S43000 ferritic stainless steel. Source: Ref 67

Despite these properties, stainless steels can be machined under the appropriate conditions. In general, more power is required to machine stainless steels than carbon steels, cutting speeds must often be lower, a positive feed must be maintained, tooling and fixtures must be rigid, chip breakers or curlers may be needed on the tools, and care must be taken to ensure good lubrication and cooling during cutting (Ref. 68). These and other practices are discussed in more detail in the sections on individual conventional


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machining techniques in the article "Machining of Stainless Steels" in Machining, Volume 16 of ASM Handbook, formerly 9th Edition Metals Handbook. Significant differences in machinability exist between different alloy systems and alloy families, including the various free-machining alloys. This section will discuss machinability both within and among the five basic families of stainless steels. Machinability of Ferritic and Martensitic Alloys. Free-machining ferritic alloys (such as S43020) and annealed, low-carbon, free-machining martensitic alloys (such as S41600) are the easiest to machine of the stainless steels (Ref 66, 69, 70, 71, 72). In fact, their machinability ratings approach and in some cases are comparable to those of certain free-machining carbon steels (Ref 69, 70, 71). The nonfree-machining lower-chromium ferritic alloys (S40500, S43000) and annealed, low-carbon, straight-chromium martensitic alloys (S40300, S41000) are also generally easier to machine than most other nonfree-machining alloys (Ref 69, 70, 71, 72). The higher-chromium ferritic alloys, such as S44600, are considered by some to be somewhat more difficult to machine than the lower-chromium alloys because of gumminess and stringy chips (Ref 73, 74). Other than the presence or lack of a free-machining additive, the machining characteristics of martensitic stainless steels are influenced by certain variables: • • • •

Hardness level Carbon content Nickel content Phase balance, that is, the percentage of free or δ-ferrite in the martensitic matrix

Increasing the hardness level for a particular alloy results in a decrease in machinability as measured by various criteria (tool life, drillability, and so on) (Ref 66, 72, 75, 76). Within certain limits, however, surface finish can be improved by machining harder material (Ref 72, 75). In the martensitic grades, machinability decreases as the carbon content increases from S41000 to S42000 to S44004 or from S41600/S41623 to S42020/S42023 to S44020/S44023. With higher carbon levels, there also tends to be a smaller difference in machinability between the corresponding free-machining and nonfree-machining versions. These effects are primarily due to the increasing quantities of abrasive chromium carbides present as carbon level increases in this series of alloys. As a further detriment to machinability, annealed hardness level increases with increasing carbon level (Ref 72, 73, 74, 75). Nickel content also influences machinability by increasing annealed hardness levels. Consequently, alloys such as S41400 and S43100 will be more difficult to machine than S41000 in the annealed condition (Ref 70, 71). Changing phase balance has been used to improve the machining characteristics of S41600. It has generally been found that increasing free or δ-ferrite content results in improved machinability, including tool life and surface finish (Ref 72, 75, 76, 77, 78, 79). The introduction of a higher ferrite content also results in a decreasing hardness capability. Machinability of Austenitic Alloys. The difficulties in machining attributed to stainless steels in general are more specifically attributable to the austenitic stainless steels (Ref 69, 71, 72, 73, 74, 79). Compared to ferritic and martensitic alloys, typical austenitic alloys have a higher work-hardening rate, a wider spread between yield and ultimate tensile strengths, and higher toughness and ductility. When machining austenitic stainless steels, particularly the nonfree-machining alloys, several factors become more pronounced: • • • •

Tools will run hotter, with more tendency to form a large built-up edge Chips will be stringier, with a tendency to tangle, making their removal difficult Chatter will be more likely if tool rigidity is inadequate or marginal Cut surfaces will be work hardened and more difficult to machine if cutting is interrupted or if the feed rate is too low

Because of these factors, the general precautions for machining stainless steels are particularly important for austenitic alloys. Although there have been differing opinions (Ref 80), a moderate amount of cold work has been regarded as beneficial to the overall machining characteristics of austenitic stainless steels (Ref 72, 74). The cold working reduces the ductility of the material, which results in cutting with a cleaner chip and less tendency for a built-up edge. This produces a better-machined surface finish but with some loss of tool life due to the higher hardness level (Ref 72). Automatic screw machine testing has shown that the effects of cold working and hardness are variable and may or may not be seen, depending on the type of alloy and the machining conditions. In such testing, tool life has been lowered by an increasing level of cold work for both free-machining (S30300) and nonfree-machining (S30400, S31600) austenitic stainless steels. This effect is shown in Fig. 33 for S30400. On the other hand, there have also been indications under different cutting conditions of an optimum level of tool life at an intermediate level of cold work (Fig. 34 ). Fig. 33 Effect of percent cold draft on machinability in a screw machine test for an enhanced-machining version of S30400. Termination is defined as a 0.075 mm (0.003 in.) increase in the diameter of the part being cut.


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Fig. 34 Effect of percent cold draft on machinability in a screw machine test for an enhanced-machining version of S30400. Termination is defined as a 0.075 mm (0.003 in.) increase in the diameter of the part being cut.

Machined surface finish can be improved by an increasing level of cold work for nonfree-machining alloys (S30400, S31600). Figure 35 shows this effect for S31600. A decreasing tendency for tool chatter with increasing cold work has also been seen for these alloys. On the other hand, the use of cold-drawn bar does not consistently benefit the machined surface finish of a free-machining alloy (S30300). Fig. 35 Effect of percent cold draft on machined surface finish in a screw machine test for an enhanced-machining version of S31600

Additions of manganese or copper can increase the machinability (Fig. 36 ) and decrease the high work-hardening rate of the


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lower-alloy austenitic stainless steels (Ref 79, 81, 82, 83, 84). Austenitic free-machining alloys that have additions of manganese and/or copper include S20300, S30310, S30330, and S30431. Although higher alloy content generally reduces the work-hardening rate, it may not necessarily benefit machinability. Highly alloyed austenitic stainless steels, such as S30900, S31000, and N08020, tend to be more difficult to machine (Ref 70, 71). Fig. 36 Effect of copper and manganese contents on machinability in a drill test for a free-machining chromium-manganese-nickel austenitic stainless steel. Source: Ref 81

Carbon and nitrogen can affect work-hardening rate and will increase the strength and hardness of austenitic stainless steels. Higher levels of either or both elements will decrease machinability (Fig. 37 ). Consequently, the high-nitrogen austenitic alloys, such as S20910 and S28200, are more difficult to machine than the standard lower-nitrogen austenitic alloys (Ref 70). Fig. 37 Effect of carbon and nitrogen contents on machinability in a tool life test for a free-machining 18Cr-9Ni-3Mn austenitic stainless steel. Source: Ref 85

Strong carbide/nitride-forming elements, including titanium and niobium, are used in stainless steels such as S32100 and S34700 to prevent grain-boundary carbide, which can reduce intergranular corrosion resistance. However, the carbide/nitride inclusions are abrasive and will increase tool wear (Fig. 38 ).


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Fig. 38 Comparison of tool wear for austenitic stainless steels with (S32100) and without (S30400) titanium carbide inclusions. Source: Ref 86

Machinability of Duplex Alloys. The machinability of duplex stainless steels is limited by their high annealed strength level. Figures 39 and 40 compare the machinability of a duplex alloy, S32950, with that of a high-nitrogen austenitic alloy, S20910, and a conventional austenitic alloy, S31600, in standard (0.004% S) and enhanced-machining (0.027% S) versions. The duplex alloy (S32950) has a hardness level comparable to that of the high-nitrogen austenitic alloy (S20910), but provides better machinability. However, it does not machine as well as either the standard or the enhanced-machining S31600 alloy. Fig. 39 Comparison of tool life for a duplex stainless steel (S32950), a high-nitrogen austenitic stainless steel (S20910), and a lower-nitrogen austenitic stainless steel (S31600). Tool life is measured as the distance traveled along a 25 mm (1.0 in.) diam bar until tool failure. Shaded areas represent distance to failure.

Fig. 40 Comparison of machinability in a drill penetration test for a duplex stainless steel (S32950), a high-nitrogen austenitic stainless steel (S20910), and a lower-nitrogen austenitic stainless steel (S31600)


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Other nitrogen-bearing duplex alloys are expected to machine similarly to S32950. No enhanced-machining versions of duplex alloys are available. Machinability of Precipitation-Hardenable Alloys. The machinability of precipitation-hardenable stainless steels depends on the type of alloy and its hardness level. Martensitic precipitation-hardened stainless steels are often machined in the solution-treated condition; therefore, only a single aging treatment is required afterward to reach the desired strength level. In this condition, the relatively high hardness limits machinability. Most of these alloys machine comparably to, or somewhat worse than, a standard austenitic alloy such as S30400. Alloy S17400 is available in enhanced-machining versions that allow machining at higher speeds with a significantly reduced tendency toward chatter. Martensitic precipitation-hardenable stainless steels can also be machined in an aged condition so that the heat treating can be avoided and closer tolerances maintained. The ease of cutting generally varies with the hardness or heat-treated condition (Table 34 ). Table 34 Relative machinability of 17-4 PH (S17400) in various heat-treated conditions Condition

Typical hardness, HRC

Improved machinability (higher cutting speed) H1150M

27

H1150

33

H1075

36

A (solution treated)

34

H1025

38

H900

44

Improved surface finish Source: Ref 87

In the annealed, austenitic condition, semiaustenitic alloys can be expected to machine with difficulty, somewhat worse than an alloy such as S30200, which has a high work-hardening rate. Alloys S35000 and S35500 can be supplied in an equalized and overtempered condition, which provides the best machinability. As with the martensitic precipitation-hardenable alloys, machining difficulties increase with aged hardness level. Austenitic precipitation-hardenable alloys, such as S66286, machine quite poorly, requiring slower cutting rates than even the highly alloyed austenitic stainless steels (Ref 70). Machining in an aged condition requires even slower speeds. General Guidelines. The characteristics of stainless steels that have a large influence on machinability include: • Relatively high tensile strength • High work-hardening rate, particularly for the austenitic alloys • High ductility These factors explain the tendency of the material to form a built-up edge on the tool during traditional machining operations. The chips removed in machining exert high pressures on the nose of the tool; these pressures, when combined with the high temperature at the chip/tool interface, cause pressure welding of portions of the chip to the tool. In addition, the low thermal conductivity of stainless steels contributes to a continuing heat buildup. Figure 41 compares the machinability ratings of selected stainless steels using AISI B1112 as the reference. The difficulties


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involved in the traditional machining of stainless steels can be minimized by observing the following points: • Because more power is generally required to machine stainless steels, equipment should be used only up to about 75% of the rating for carbon steels • To avoid chatter, tooling and fixtures must be as rigid as possible. Overhang or protrusion of either the workpiece or the tool must be minimized. This applies to turning tools, drills, reamers, and so on • To avoid glazed, work-hardened surfaces, particularly with austenitic alloys, a positive feed must be maintained. In some cases, increasing the feed and reducing the speed may be necessary. Dwelling, interrupted cuts, or a succession of thin cuts should be avoided • Lower cutting speeds may be necessary, particularly for nonfree-machining austenitic alloys, precipitation-hardenable stainless steels, or higher-hardness martensitic alloys. Excessive cutting speeds result in tool wear or tool failure and shutdown for tool regrinding or replacement. Slower speeds with longer tool life are often the answer to higher output and lower costs • Tools, both high-speed steel and carbides, must be kept sharp, with a fine finish to minimize friction with the chip. A sharp cutting edge produces the best surface finish and provides the longest tool life. To produce the best cutting edge on high-speed steel tools, 60-grit roughing should be followed by 120- and 150-grit finishing. Honing produces an even finer finish • Cutting fluids must be selected or modified to provide proper lubrication and heat removal. Fluids must be carefully directed to the cutting area at a sufficient flow rate to prevent overheating Fig. 41 General comparison of machinability of stainless steels compared with AISI B1112. Rating based on 100% for AISI B1112 using high-speed steel tools. Source: Ref 88


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Weldability The metallurgical features of each group generally determines the weldability characteristics of the steels in that group. The weldability of martensitic stainless steels is greatly affected by hardenability that can result in cold cracking. Welded joints in ferritic stainless steel have low ductility as a result of grain coarsening that is related to the absence of allotropic (phase) transformation. The weldability of austenitic stainless steels is governed by their susceptibility to hot cracking, as is the case with other single-phase alloys with a face-centered cubic (fcc) crystal structure. With the precipitation-hardening stainless steels, weldability is related to the mechanisms associated with the transformation (hardening) reactions (Ref 89). Stainless steels can be joined by most welding processes, but with some restrictions. In general, those steels that contain aluminum or titanium, or both, can be arc welded only with the gas-shielded processes. The weld joint efficiency depends upon the ability of the welding process and procedures to produce nearly uniform mechanical properties in the weld metal, heat-affected zone, and base metal in the as-welded or postweld heat-treated condition. These properties can vary considerably with ferritic, martensitic, and precipitation hardening steels (Ref 89). Weldability and various suitability-for-service conditions, including temperature, pressure, creep, impact, and corrosion environments (see the article "Corrosion of Weldments" in Corrosion, Volume 13 of ASM Handbook, formerly 9th Edition Metals Handbook), require careful evaluation because of the complex metallurgical aspects of stainless steels. Shielded metal arc (SMAW), submerged arc (SAW), gas metal arc (GMAW), gas tungsten arc (GTAW), and plasma arc welding (PAW) are used extensively for joining stainless steels. Flux cored arc welding (FCAW) is also used, but to a lesser extent. The remainder of the article discusses the weldability of the various grades and the suitability of arc welding processes for specific conditions and requirements.


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Austenitic Stainless Steels. Differences in composition among the standard austenitic stainless steels affect weldability and performance in service. For example, types 302, 304, and 304L differ primarily in carbon content, and consequently there is a difference in the amount of carbide precipitation that can occur in the heat-affected zone (HAZ) after the heating and cooling cycle encountered in welding. Types 303 and 303Se contain 0.20% P (max) plus 0.15% Se or S for free machining. These elements are detrimental to weldability and can cause severe hot cracking in the weld metal. Types 316 and 317 contain molybdenum for increased corrosion resistance and higher creep strength at elevated temperatures. However, unless controlled by extralow carbon content, as in type 316L, carbide precipitation occurs in the HAZ during welding. Types 318, 321, 347, and 348 are stabilized with titanium, or niobium-plus-tantalum, to prevent the intergranular precipitation of chromium carbides when the steels are heated to a temperature in the sensitizing range, as during welding. Welding Characteristics. The austenitic stainless steels, except for the free-machining grades, are the easiest to weld and produce welded joints that are characterized by a high degree of toughness, even in the as-welded condition. Serviceable joints can be readily produced if the composition and the physical and mechanical properties are well suited to the welding process and condition. The heat of welding, contamination, carbide precipitation, cracking, and porosity must be considered before, during, and after welding stainless steels. Heat of Welding. Excessive heat input may result in weld cracking, loss of corrosion resistance, warping, and undesirable changes in mechanical properties. Welds in stainless steels generally require 20 to 30% less heat input than welds in carbon grades because stainless steels have lower thermal conductivity and higher electrical resistance. Because of low thermal conductivity, heat remains near the weld, so that more heat is available to melt the material, which may produce detrimental results. Excessive heat produces large thermal gradients across the joint, which can cause distortion. Because heat dissipates slowly in stainless steel, it may lower corrosion resistance and change strength. These effects can be minimized with chill bars and less heat input. Weld metal cracking, another problem resulting from excessive welding heat, is discussed later in this article. The high electrical resistivity of stainless steel makes it suitable for welding with low heat inputs. With reduced heat, good penetration and fusion result because low thermal conductivity retains heat in the weld area. Comparative electrical resistivities are 25 to 50 µΩ · in. for carbon steel and 175 to 200 µΩ · in. for austenitic grades. Heat input can be reduced by using low amperages, low voltages (short arc lengths), high travel speeds, and stringer beads. With GMAW and GTAW processes, heat input can be affected by the type of shielding gas. Argon produces a cool, stable arc, while helium produces a hot arc that is somewhat unstable. For manual processes, pure argon is generally best. When working with automatic welding equipment that offers good control of amperage, voltage, and travel speed, mixed gases can be used without risking damage from high heat. Finally, pulsed arc welding techniques can be used to lower heat input. Weld Contamination. Contaminants not only hinder successful welding, but may prevent an apparently sound weld from functioning satisfactorily. A contaminated weld has inferior corrosion resistance and strength, and the weldment may fail prematurely. The stainless steel itself may also contain the contaminant. Free-machining stainless steels frequently contain sulfur or selenium. Both elements can make the steel unweldable. Similarly, high concentrations of carbon in high-strength stainless steel can inhibit weld serviceability. External sources of contamination include carbon, nitrogen, oxygen, iron, and water. Carbon is often picked up from shop dirt, grease, forming lubricants, paint, marking materials, and tools; consequently, steel parts should be cleaned before welding and during welding. Otherwise, carbon contamination can cause welds to crack, change the mechanical properties, and lower the corrosion resistance in weld areas. Although iron contamination generally does not affect weldability, it can lower serviceability. Flakes of iron on surfaces rust, thereby speeding localized corrosion. The welder may unknowingly cause the contamination by grinding stainless steel with a wheel previously used on carbon steel. Clean Al2O3 grinding wheels, preferably those not used for grinding other alloys, should be used. One of the most troublesome types of contamination is stainless steel surface contamination by copper, bronze, lead, or zinc from hammers, hold-down fingers on seamers, or tools used in fabrication. Small amounts of these materials on the surface of the stainless steel can lead to cracking in the high-temperature HAZ of the weld. This type of cracking generally occurs in the HAZ, where the contaminant attacks the grain boundaries. Effect of Carbide Precipitation on Corrosion Resistance of Welded Joints. The precipitation of intergranular chromium carbides is accelerated by an increase in temperature within the sensitizing range and by an increase in time at temperature. When intergranular chromium carbides are precipitated at welded joints, resistance to intergranular corrosion and stress corrosion markedly decreases. The decrease in corrosion resistance is attributed to the presence of the chromium-rich carbides at the grain boundaries and the depletion of chromium in the adjacent matrix material. Although intergranular carbide precipitation generally occurs between 425 and 870 °C (800 and 1600 °F), sensitization is restricted to a narrow range by the fairly rapid heating and cooling that usually occur in welding. The narrower range varies with time at temperature and steel composition, but is approximately 650 to 870 °C (1200 to 1600 °F). The base metal immediately adjacent to the weld is annealed or solution treated by the heat of welding and, because it generally is cooled rapidly enough to hold the dissolved carbides in solution, this zone usually exhibits normal resistance to corrosive attack. A short distance from the weld, about 3.2 mm (1=8in.) (the distance depending on the thermal cycle and material thickness), there is a narrow zone in which lower heating and cooling rates prevail. In this HAZ, intergranular precipitation of chromium carbides is most likely to take place. Harmful carbide precipitation can be overcome or prevented by the use of: • Postweld solution annealing


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• Extralow-carbon, that is, 0.03% C (max), alloy • Stabilized alloy containing preferential carbide-forming elements, such as niobium-plus-tantalum or titanium Solution annealing puts carbides back into solution and restores normal corrosion resistance, but is generally inconvenient. The solution-annealing temperature range is very high, 1040 °C (1900 °F) minimum, and unless stainless steel is protected from air at these temperatures, it oxidizes rapidly, forming adherent oxide scale. Thin sections, unless adequately supported, may sag or be severely distorted at these temperatures or during rapid cooling from them. Rapid cooling in solution annealing may present other problems. Water quenching, although effective, is seldom feasible except for small workpieces of simple shape. Unless adequate safeguards are available, water quenching of large workpieces from the solution-annealing temperature is hazardous. Often solution annealing is impractical because the workpiece is too large for available furnace and cooling facilities. Extralow-carbon stainless steels (types 304L and 316L) are resistant to carbide precipitation in the 425 to 870 °C (800 to 1600 °F) range and can thus undergo normal welding without reduction in corrosion resistance. Carbides precipitate in significant quantities when extralow-carbon steels are heated and held in the sensitizing temperature range for an extended period, as in service. These steels are generally recommended for use below 425 °C (800 °F). Stabilized Steels. Compared with the extralow-carbon steels, the stabilized steels exhibit higher strength at elevated temperature. For service in a corrosive environment in the sensitizing temperature range of 425 to 870 °C (800 to 1600 °F), an austenitic steel stabilized with niobium-plus-tantalum or titanium is needed. The filler metal used for welding should also be of a stabilized composition. Because an inert shielding gas is used, GTAW and GMAW are suitable for titanium-stabilized steel without oxidizing the titanium. Under certain conditions, stabilized stainless steel weldments are susceptible to sensitization, which occurs in narrow zones of the base metal immediately adjacent to the line of weld fusion. During welding, stabilized carbides are dissolved and, as a result of rapid cooling, are retained in solution. Subsequent reheating to about 650 °C (1200 °F) results in preferential precipitation of chromium carbides in a narrow zone that exhibits less than normal corrosion resistance. Microfissuring in Welded Joints. Interdendritic cracking in the weld area that occurs before the weld cools to room temperature is known as hot cracking or microfissuring. The occurrence of microfissuring is related to the: • • • • •

Microstructure of the weld metal as solidified Composition of the weld metal, especially the content of certain residual or trace elements Amounts of stress developed in the weld as it cools Ductility of the weld metal at high temperatures Presence of notches

Susceptibility to microfissuring is highly dependent on the microstructure of the weld metal. Weld metal with a wholly austenitic microstructure is considerably more susceptible to microfissuring than weld metal with a duplex structure of δ-ferrite in austenite. The content of alloying elements and residual elements strongly influences the susceptibility of fully austenitic stainless steel weld metal to microfissuring. Susceptibility can be reduced by a small increase in carbon or nitrogen content or by a substantial increase in manganese content. Residual or trace elements that contribute to microfissuring are boron, phosphorus, sulfur, selenium, silicon, niobium, and tantalum. The amount of stress imposed on austenitic stainless weld metal as it cools from the solidus down to about 980 °C (1800 °F) should be minimized. In this temperature range, the weld metal is most susceptible to microfissuring, and if the level of stress is high, the fissures propagate to form visible cracks. Peening is not an effective method of preventing this type of cracking because it can seldom be applied early enough to reduce stress buildup. Prevention of Microfissuring. To obtain duplex-structured weld metal that has a controlled ferrite content of at least 3 to 5 ferrite number (FN), a filler metal of suitable composition is selected. The ferrite number is a magnetically determined scale of ferrite measurement. The Welding Research Council Advisory Subcommittee on Welding Stainless Steels determined that the ferrite number of a weld metal, at least from 0 to 6 FN, approximates the average value of percent ferrite assigned by laboratories applying metallographic measurements of ferrite to a given weld metal. Microfissuring can be prevented or minimized by the proper control of ferrite in the weld metal. Wide use has been made of the Schaeffler diagram (Fig. 42 ) to determine the approximate amount of ferrite that will be obtained in the austenitic weld metal of a given composition. Point X indicates the equivalent composition of a type 318 (316Cb) weld deposit containing 0.07C-1.55Mn-0.57Si-18.02Cr-11.87Ni-2.16Mo-0.80Nb. To determine the chromium and nickel equivalents, each percentage was multiplied by the potency factor indicated for the respective element along the axes of the diagram. When these were plotted as point X, the constitution of the weld was indicated as austenite plus 0 to 5% ferrite. Magnetic analysis of an actual sample revealed an average 2 FN. For austenite-plus-ferrite structures, the diagram predicts ferrite within 4% for stainless steel types 308, 309, 309Cb, 310, 312, 316, 317, and 318. Actual measurements of ferrite content can be made conveniently with the aid of a magnetic analysis device. American Welding Society (AWS) A4.2-74, "Standard Procedures for Calibrating Magnetic Instruments to Measure the Delta Ferrite Content of Austenitic Stainless Steel Weld Metal," and A5.4-81, "Specification for Covered Corrosion-Resisting Chromium and Chromium-Nickel Steel Welding Electrodes," discuss this measurement. Fig. 42 Schaeffler diagram. See Fig. 30 for revised Schaeffler diagram.


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Because many heats of austenitic stainless steel contain appreciable amounts of nitrogen (a very strong austenitizer), a revised constitution diagram for austenitic stainless steel weld metal has been developed to include nitrogen in the nickel equivalent (Fig. 30 ). Compared to Fig. 42 , Fig. 30 is modified in shape and slope to improve the accuracy of ferrite estimation for types 309, 309Cb, 316, 316L, 317, 317L, and 318. In addition, ferrite calculation for types 308 and 347 weld metal is improved on samples with either high or low nitrogen content. For use in the diagram, actual nitrogen content is preferred. If it is not available, the following nitrogen value shall be used: 0.12% for self-shielding flux cored electrode GMAW welds, 0.08% for other GMAW welds, and 0.06% for welds of other processes. When the weld metal must be wholly austenitic (that is, when the metal must be nonmagnetic or when specific corrosive environments that selectively attack δ-ferrite will be encountered), the content of crack-promoting residual elements must be stringently controlled, and the composition of the weld metal must be adjusted to increase crack resistance. Crack resistance can be increased by modifying the carbon, manganese, sulfur, phosphorus, silicon, and nitrogen contents of the weld metal. However, even with optimum compositions and the most favorable welding procedures, wholly austenitic weld deposits are more crack sensitive than those of a duplex structure. Underbead cracking can occur in the heat-affected zones adjacent to welds in austenitic stainless steel, particularly when the weld zone is heavily retrained or the section thickness is greater than 19 mm (3=4in.). Such cracking is most common in type 347 because of the strain-induced precipitation of niobium carbides, but it has been reported in other types as well. Weld restraint is the most important factor in the control of underbead cracking. Preheating is of little value in preventing cracking of austenitic stainless steel and can cause other problems, such as increased carbide precipitation. Selection of Filler Metals. Electrodes and welding rods suitable for use as filler metal in the welding of austenitic stainless steels are shown in Table 35 . These filler metals, with AWS standard composition specifications, are for GMAW, SAW, and SMAW. The selection of filler metals for welding austenitic stainless steels requires consideration of the microstructural constituents of the as-deposited weld metal. Ultimately, these microstructural constituents determine the mechanical properties, crack sensitivity, and corrosion resistance of the weld. The constituents of principal concern are austenite, δ-ferrite, and precipitated carbides. Table 35 Suggested filler metals for stainless steels Recommended filler metal(a) AISI No.

First choice

Second choice

Remarks

Chromium-nickel-manganese austenitic nonhardenable 201

308

308L

Substitute for 301

202

308

308L

Substitute for 302

301

308

308L

302

308

308L

302B

308

309

303

...

...

Free-machining stainless steel: welding not recommended⎯312

303Se

...

...

Free-machining stainless steel: welding not


High silicon

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01 Sep 2005

recommended⎯312 Chromium-nickel austenitic nonhardenable 304

308

308L

308L

347

305

308

...

308

308

...

304L

309

309

...

309S

309

...

310

310

...

310S

310

...

314

310

...

316

316

309Cb

316L

309Cb

317

317

309Cb

321

347

308L

347

347

308L

348

347

...

316L

Extralow carbon

Low carbon Low carbon

Extralow carbon

Difficult to weld in heavy sections

Chromium martensitic hardenable 403

410

...

410

410

430

414

410

...

416

410

...

416Se

410-15 should be used

...

...

Free-machining: welding not recommended

420

410

...

High carbon

431

430

...

440A

...

...

High carbon: welding not recommended

440B

...

...

High carbon: welding not recommended

440C

...

...

High carbon; welding not recommended

Chromium ferritic nonhardenable 405

410

405Cb

430F

...

...

Free machining: welding not recommended

430FSe

...

...

Free machining: welding not recommended

446 309 310 (a) Use E, electrode, or R, welding rod, prefix. Source: Ref 90

Some filler metals, such as types 309Cb and 310, invariably deposit a fully austenitic weld metal. In these alloys, the ratio of ferrite formers to austenite formers cannot, within permissible limits, be raised high enough to produce any δ-ferrite in the austenite. Consequently, when these filler metals are applied to restrained joints or to base metals containing additions of elements such as phosphorus, sulfur, selenium, and silicon, only those procedures proved suitable by experience should be used. The compositions of most filler metals are adjusted by the manufacturers to produce weld deposits that have ferrite-containing microstructures. Thus, ferrite-forming elements, such as chromium and molybdenum, are maintained at the high side of their allowable ranges, and austenite-forming elements, such as nickel, are kept low. The amount of ferrite in the structure of the weld metal depends on the ratio or balance of these elements. At least 3 or 4 FN δ-ferrite is needed in the as-deposited weld metal for effective suppression of hot cracking. With the proper techniques, however, types 316 and 316L can be welded with as little as 0.5 FN. Ferrite-containing weld metal may have certain disadvantages in a welded austenitic stainless steel. Ferrite is ferromagnetic, and the increased magnetic permeability of the weld metal may be objectionable in applications that require nonmagnetic properties. When exposed to service at elevated temperature, the ferrite in some weld metals may transform to σ phase and adversely affect mechanical properties and corrosion resistance, a problem that has been encountered in power plant applications. When a joint is arc welded without the addition of filler metal, the structure of the weld metal is determined by the composition of the base metal. Sometimes this leads to unfavorable results, because wrought base metals may not have the compositional limits required for good weld metal. Preheating. In general, no benefit is derived from preheating austenitic stainless steels. In some applications, preheating can increase carbide precipitation, cause shape distortion of the workpiece, or increase hot-cracking tendencies.


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Postweld Stress Relieving. Although the effects of residual stress from welding on the properties of austenitic stainless steels are limited in comparison to the effects of cold working, residual stress may significantly affect the mechanical properties. Because the effective yield strength varies from point to point, the application of further stresses at later stages of fabrication can cause excessive distortion and even premature failure. Nonuniform heating, which relieves some local residual stress, may also contribute to distortion. For these reasons, stress relieving may be required to ensure dimensional stability. Stress relieving can be performed over a wide range of temperatures, depending on the amount of relaxation required. Time at temperature ranges from about 1 h per inch of section thickness at temperatures above 650 °C (1200 °F) to 4 h per inch of section thickness at temperatures below 650 °C (1200 °F). Because of the high coefficient of expansion and the low thermal conductivity of austenitic stainless steels, cooling from the stress-relieving temperature must be slow. The stress-relieving temperature selected must be compatible with the extent of carbide precipitation acceptable and with the corrosion resistance desired. Nonstabilized stainless steels cannot be stress relieved in the sensitizing temperature range without sacrifice of corrosion resistance. Extralow-carbon stainless steels are affected much less, because carbide precipitation in these steels is sluggish. Stabilized stainless steels exhibit minimal chromium carbide precipitation tendencies. For austenitic stainless steels, the estimated percentages of residual stress relieved at various temperatures, for the times previously noted, are: Temperature °C

°F

Stress relief, %

845−900

1550−1650

85

540−650

1000−1200

35

Ferritic Stainless Steels. The ferritic stainless steels are generally less weldable than the austenitic stainless steels and produce welded joints having lower toughness because of grain coarsening that occurs at the high welding temperatures. The standard ferritic stainless steels are: • Type 446 (25% Cr) • Types 430, 430F, and 430F-Se (17% Cr) • Types 405 and 409 (13% Cr) Type 409 is ferritic because it has a low carbon content (0.08% max) and a minimum titanium content equal to six times the carbon content. Type 405, which also contains only 0.08% C (max), contains an average of 0.20% Al, which promotes ferrite formation. There are a number of ferritic stainless steels that contain very low amounts of carbon and nitrogen. The low-interstitial ferritic stainless steels (including the 26Cr-1Mo and 29Cr-4Mo alloys) can be welded so that they do not lose any of the base metal ductility in the weld area, while retaining grain growth. The key to the successful welding of these steels is to prevent any carbon, nitrogen, or oxygen contamination during welding. Thus, the part and filler material must be clean before welding, and both the molten weld metal and the hot weld-area metal must be fully shielded from the atmosphere. All moisture must be excluded from the weld area before and during welding. Effect of Welding Heat on Ductility and Grain Size. Although most ferritic stainless steels have compositions that ensure a ductile ferritic structure at room temperature, variations in composition within the standard compositional limits can result in the formation of small amounts of austenite during heating to elevated temperature. Upon cooling, the austenite transforms to martensite, resulting in a duplex structure of ferrite and a small amount of martensite. The martensite reduces both the ductility and toughness of the steel. Annealing transforms the martensite and restores normal ferritic properties, but annealing increases costs and can result in an excessive amount of distortion, particularly in parts that were previously formed by a cold-working process. All ferritic stainless steel mill products are normally annealed at the mill to transform any martensite that may be present to a softer structure of ferrite and carbides. In this condition, the steel can be readily cold formed. Only when the steel is heated near or above the transformation temperature (approximately 870 °C, or 1600 °F), as during welding, does the risk of austenite formation and subsequent transformation to martensite arise. In addition, heating to temperatures above 955 °C (1750 °F) results in enlargement of the ferrite grain size, which also reduces the ductility and toughness of the steel. Although martensite can be eliminated by annealing, coarsened ferrite grains remain unaffected. Because martensite responds to annealing and inhibits ferrite grain growth, some applications may benefit from martensite formation, provided that the workpiece can be annealed after welding. In a 17% Cr steel, martensite formation is promoted by lowering the chromium content to 15 to 16%. When this practice is adopted, it is usually necessary to preheat before welding or to select a steel of lower carbon content to guard against cracking in the HAZ. When postweld annealing is not feasible, the ductility of the welded joint can be controlled by the selection of a stainless steel base metal containing substantial amount of strong ferrite-former, such as aluminum, niobium, or titanium. One such steel, which has the commercial designation of type 430Ti, is a 17% Cr steel containing 0.12% C (max) and a minimum titanium content equal to six times the carbon content. The metallurgical functions of the titanium in this steel are to form stable titanium carbides and to promote the formation of ferrite. When a completely ferritic steel is welded, no martensite is formed in the HAZ, although some grain coarsening may occur. Grain coarsening can be controlled, to some extent, by minimizing heat input during welding


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and by avoiding slow cooling from the weld temperature. Effect of Temperature on Notch Toughness. For the 17% Cr steels, the temperature range is just above room temperature for transition from a tough shear-type fracture at the higher temperature to a brittle, cleavage-type fracture at the lower temperature, upon impact at a notch. Under impact loading at room temperature and below, these steels are notch sensitive, and impact test values of less than 20 J (15 ft · lbf) are usual. At 95 to 120 °C (200 to 250 °F), the impact test values increase to approximately 40 to 70 J (30 to 50 ft · lbf). This relationship between notch toughness and temperature is important in the selection of joint design and welding conditions for ferritic stainless steel. In service, a weldment designed primarily to withstand static load may be subjected to accidental or unforeseen impact loading. Furthermore, a weldment with low notch toughness may not withstand an appreciable number of cyclical stresses even under a low rate of loading. Because multiaxial residual stresses are often developed during welding (especially when welding heavy sections), notches and points of stress concentration that might cause failure in service must be avoided whenever possible. Preheating before welding is often useful in preventing cracking during welding. Preheating. The recommended preheating temperature range for ferritic stainless steels is 150 to 230 °C (300 to 450 °F). The need for preheating is determined largely by the composition, mechanical properties, and thickness of the steel being welded. Steels less than 6.4 mm (1=4in.) thick are much less likely to crack during welding than those of greater thickness. The type of joint, joint location, restraints imposed by clamping and jigging, welding process, and rate of cooling from the welding temperature can also affect weld cracking. Postweld Annealing. The temperature range for postheating or postweld annealing of ferritic stainless steels is 790 to 845 °C (1450 to 1550 °F), which is safely below the temperatures for austenitic formation and grain coarsening. Annealing transforms a mixed structure to a wholly ferritic structure and restores the mechanical properties and corrosion resistance that may have been adversely encountered in welding. Thus, except for its inability to refine coarsened ferrite grains, annealing is generally beneficial. Annealing has two major disadvantages: the time and cost of the treatment, and the need to prevent the formation of the oxide scale or to remove it if it is already present. Annealing may also require the use of elaborate fixturing to prevent sagging or distortion of the weldment. Cooling ferritic stainless steel from the annealing temperature may be done by air or water quenching. To minimize distortion from handling, weldments are often allowed to cool to about 595 °C (1100 °F) before they are removed from the furnace. Slow cooling through the temperature range of 565 to 400 °C (1050 to 750 °F) must be avoided because it produces brittleness in the steel. Susceptibility to this type of embrittlement, known as 475 °C (885 °F) embrittlement, normally increases as chromium content increases. Heavy sections may require forced cooling or a spray quench to bring them safely through this embrittlement range. Selection of Filler Metal. As shown in Table 35 , both ferritic and austenitic stainless steel filler metals are used in the arc welding of ferritic stainless steel. Ferritic stainless steel filler metals offer the advantages of having the same color and appearance, the same coefficient of thermal expansion, and essentially the same corrosion resistance as the base metal. However, austenitic stainless steel filler metal are often used to obtain more ductile weld metal in the as-welded condition. Although austenitic stainless steel weld metal does not prevent grain growth or martensite formation in the HAZ, the ductility of austenitic weld metal improves the ductility of the welded joint. The selection of austenitic stainless steel filler metal, however, should be carefully related to the specific application to determine whether differences in color or in the physical corrosion and mechanical properties of the weld metal and the base metal cause difficulty. For weldments that are to be annealed after welding, the use of austenitic filler metal can introduce several problems. The normal range of annealing temperature for ferritic stainless steels falls within the sensitizing temperature range for austenitic steels. Consequently, unless the austenitic weld metal is of extralow-carbon content or is stabilized with niobium or titanium, its corrosion resistance may be seriously impaired. If the annealing treatment is intended to relieve residual stress in the weldment, it cannot be fully effective because of the difference in the coefficients of thermal expansion of the weld metal and the base metal. Corrosion Resistance. Ferritic stainless steels usually exhibit less corrosion resistance than austenitic stainless steels. Any condition of the welded joint that might therefore impair corrosion resistance must be avoided. The presence of martensite or the precipitation of σ phase at the grain boundaries can cause severe intergranular corrosion in the HAZ. Completely ferritic steels, such as types 430Ti and 446, display little or no susceptibility to intergranular attack at the weld joint. Annealing of any welded ferritic steel eliminates the unfavorable structural conditions that promote corrosive attack. Martensitic Stainless Steels. The standard martensitic stainless steels are types 403, 410, 414, 416, 416Se, 420, 431, 440A, 440B, and 440C. These steels derive their corrosion resistance from chromium, which they contain in proportions ranging from 11.5 to 18%. Martensitic stainless steels are the most difficult stainless steels to weld because they are chemically balanced to become harder, stronger, and less ductile through thermal treatment. These same metallurgical changes occur from the heat of welding. As a result of welding, these changes are restricted to the weld area and are not uniform over the entire section. The nonuniform metallurgical condition of the part makes it susceptible to cracking when subjected to the high stresses from welding. Increasing carbon content in martensitic stainless steels generally results in increased hardness and reduced ductility. Thus, the three type 440 stainless steels are seldom considered for applications that require welding, and filler metals of the type 440 compositions are not readily available. Modifications of the standard martensitic steels contain additions of elements such as nickel, molybdenum, vanadium, and tungsten, primarily to raise the allowable service temperature above the 595 °C (1100 °F) limit for the standard steels. When these elements are added, carbon content is increased, and the problem of avoiding cracking in the hardened HAZ of weldments becomes more serious.


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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. The hardness of the HAZ depends primarily on the carbon content of the base metal. As hardness increases, toughness decreases, and the zone becomes more susceptible to cracking. Preheating and control of the interpass temperature are the most effective means of avoiding cracking. Postweld heat treatment is required to obtain optimum properties. Preheating and Postweld Heat Treating. The usual preheating temperature range of martensitic steels is 205 to 315 °C (400 to 600 °F). The carbon content of the steel is the most important factor in determining whether preheating is necessary. On the basis of carbon content alone, a steel containing not more than 0.10% C seldom requires preheating, and one with more than 0.10% C requires preheating to prevent cracking. Other factors that determine the need for preheating are the mass of the joint, degree of restraint, presence of a notch effect, and composition of the filler metal. Correlations of preheating and postweld heat-treating practice with carbon contents and welding characteristics of martensitic stainless steels can be used: • Carbon below 0.10%: Neither preheating nor postweld annealing is generally required; steels which carbon contents this low are not standard • Carbon 0.10 to 0.20%: Preheat to 260 °C (500 °F); weld at this temperature; cool slowly • Carbon over 0.20 to 0.50%: Preheat to 260 °C (500 °F); weld at this temperature; anneal • Carbon over 0.50%: Preheat to 260 °C (500 °F); weld with high heat input; anneal If the weldment is to be hardened and tempered immediately after welding, annealing may be omitted. Otherwise, the weldment should be annealed immediately after welding, without cooling to room temperature. The functions of a postweld heat treatment are: • To temper or anneal the weld metal and HAZ to optimize hardness, toughness, and strength for the intended application • To decrease residual stresses associated with welding Postweld heat treatments normally used for martensitic stainless steels are subcritical annealing and full annealing (Ref 89). The necessity for a postweld heat treatment depends on the composition of the steel, the filler metal, and the service requirements. Full annealing transforms a multiple-phase weld zone to a wholly ferritic structure. This annealing procedure requires proper control of the complete thermal cycle. It should not be used unless maximum softness is required because of the formation of coarse carbides in the microstructure that take longer to dissolve at the austenitizing temperature. Typical postweld annealing temperatures are given in Table 36 . Table 36 Annealing treatments for martensitic stainless steels Subcritical annealing temperature range, °C ( °F)(a)

Full annealing temperature range, °C ( °F)(b)

403, 410, 416

650−760 (1200−1400)

830−885 (1525−1625)

414

650−730 (1200−1350)

Not recommended

420

675−760 (1250−1400)

830−885 (1525−1625)

431

620−705 (1150−1300)

Not recommended

440A, 440B, 440C

675−760 (1250−1400)

845−900 (1550−1650)

CA-6NM

595−620 (1100−1150)

790−815 (1450−1500)

Type

CA-15, CA-40

620−650 845−900 (1150−1200) (1550−1650) (a) Air cool from temperature; lowest hardness is obtained by heating near the top of the range. (b) Furnace cool to 595 °C (1100 °F); weldment can then be air-cooled. Source: Ref 89

Precipitation-Hardening Stainless Steels. Steels are divided into three groups on the dual basis of characteristic alloying additions, particularly the elements added to promote precipitation hardening, and the matrix structures of the steels in the solution-annealed and aged condition. Because differences among the steels have a direct bearing on the behavior of the steels in heat treatment and welding, the metallurgical characteristics of each are considered separately. Martensitic Precipitation-Hardening Steels. These steels have a predominantly austenitic structure at the solution-annealing


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temperature of approximately 1040 to 1065 °C (1900 to 1950 °F), but they undergo an austenite-to-martensite transformation when cooled to room temperature. The temperature at which martensite starts to form from austenite upon cooling, Ms, is usually in the range of 95 to 150 °C (200 to 300 °F). When martensite is reheated to 480 to 595 °C (900 to 1100 °F), precipitation hardening and strengthening occur, promoted by the presence of one or more alloying additions. Molybdenum, copper, titanium, niobium, and aluminum (and their compounds) are dissolved during annealing and retained in solid solution by rapid cooling, producing precipitate (usually submicroscopic particles) that increases both the strength and the hardness of the martensitic matrix. The compositional balance in the martensitic precipitation-hardening (PH) steels is critical, because relatively slight variations can lead to the formation of excessive amounts of δ-ferrite during solution annealing. If the austenite is too stable, large amounts of austenite can also be retained at room temperature after solution annealing. Either of these two conditions prevents full hardening during aging. Carbon and nitrogen contents can significantly affect this balance. Increased carbon or nitrogen may result in contamination. Typical sources of these contaminants are shop dirt and the atmosphere. These steels can be readily welded. The welding procedures resemble those ordinarily used for the 300-series stainless steels, despite differences in composition and structure between the two classes. The formation of martensite, which occurs during cooling from elevated temperatures, as in welding, does not result in full hardening. These steels are not sensitive to cracking and do not require preheating. The selection of filler metal depends on the properties required for the welded joint (Table 37 ). If strength comparable to that of the base metal is not required at the welded joint, a tough 300-series stainless steel filler metal may be adequate. When a weld having mechanical properties comparable to those of the hardened base metal is desired, the filler metal must be of comparable composition, although slight modifications are permissible to obtain better weldability. Table 37 Recommended filler metals for welding precipitation-hardening stainless steels Type

UNS designation

Covered electrodes

Bare welding wire

Dissimilar PH stainless steels

Martensitic 17-4 PH

S 17400

AMS 5827B (17-4 PH) or E308

AMS 5826 (17-4 PH) or ER308

E309 or ER309, E309Cb or ER309Cb

15-5 PH

S 15500

AMS 5827B (17-4 PH) or E308

AMS 5826 (17-4) or ER308

E309 or ER309, E309Cb or ER309Cb

17-7 PH

S 17700

AMS 5827B (17-4 PH), E308, AMS 5824A (17-7 PH) or E309

E310 or ER 310, ENiCrFe-2, or ERNiCr-3

PH 15-7 Mo

S 15700

E308 or E309

AMS 5812C (PH 15-7 Mo)

E309 or ER309, E310 or ER310

AM350

S 35000

AMS 5775A (AM350)

AMS 5774B (AM 350)

E308 or ER308, E309 or ER309

AM355 S 35500 AMS 5781A (AM355) See AWS A5.11−85. See AWS A5.14−89. Source: Ref 89

AMS 5780A (AM 355)

E308 or ER308, E309 or ER309

Semiaustenitic

When welds are deposited in a single pass, the weld metal and the HAZ usually respond uniformly to a postweld precipitation-hardening heat treatment. There is seldom any significant variation in hardness across the joint. Multiple-pass weld, however, exhibit less uniform response to the same heat treatment because successive applications of heat during welding result in marked variations in the structure of weld metal, HAZ, and base metal. Annealing eliminates these variations and provides a more uniform microstructure capable of responding uniformly to precipitation hardening. Semiaustenitic Precipitation-Hardening Steels. Unlike martensitic PH steels, semi-austenitic PH steels are soft enough in the annealed condition to permit cold working. When cooled rapidly from the annealing temperature to room temperature, they retain their austenitic structure, which displays good toughness and ductility in cold-forming operations. The M s temperatures for these steels are well below room temperature, but they vary depending on composition and annealing temperature. To obtain hardening and strengthening, the austenitic structure must be transformed to an essentially martensitic one. This can be accomplished by treating the steel before subjecting it to the precipitation-hardening heat treatment by: • Heating the steel in the range of 650 to 870 °C (1200 to 1600 °F) to precipitate carbides and other compounds, thereby depleting the matrix of enough austenite-stabilizing elements to allow transformation of austenite to martensite when the steel is cooled to room temperature • Refrigerating the steel to a temperature well below the Ms point (−75 °C, or −100 °F, for example) • Cold working the steel enough so that the austenite transforms to martensite After transformation to martensite, the semiaustenitic PH steels, like the martensitic PH steels, respond to precipitation hardening in the temperature range of 455 to 595 °C (850 to 1100 °F). Whether a precipitate forms or a tempering reaction takes place depends on the steel composition. The Ms temperature of the semiaustenitic PH steels is controlled by the solution-annealing temperature, as well as by


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composition. For example, when AM-350 steel is solution annealed at temperatures below 925 °C (1700 °F), incomplete carbide solution raises the Ms temperature above room temperature. On the other hand, when the solution-annealing temperature is raised above 925 °C (1700 °F), the Ms temperature drops precipitously. In practice, the solution-annealing temperature is not permitted to exceed about 1050 °C (1925 °F), because higher temperatures promote the formation of δ-ferrite. The semiaustenitic PH steels are normally welded in the annealed condition. The tough austenitic structure imparts welding characteristics similar to those of 300-series stainless steels. The semiaustenitic PH steels are not susceptible to cracking when welded, even when welded after transformation to martensite, because the low-carbon martensite developed is not of high hardness or low ductility. Also, cold cracking does not occur in the base metal adjacent to the weld because the HAZ is austenitized during welding and remains substantially austenitic as the joint cools to room temperature. The choice of filler metal depends largely on the weld properties desired. The filler metal can be an alloy of precipitation-hardening composition capable of developing mechanical properties comparable to those of the base metal (Table 37 ). If high strength is not requisite, the filler metal can be a 300-series austenitic stainless steel. When these steels are welded in the annealed condition, certain microstructural relations are generally obtained as a result of relatively rapid heating and cooling at the joint: • Weld metal contains small amounts of ferrite in an essentially austenitic matrix; hardness is approximately 90 HRB • Base metal immediately adjacent to the weld displays high-temperature annealed (austenitic) structure; hardness is approximately 90 HRB • Base metal in the narrow zone just beyond the annealed zone next to the weld is hardened slightly; hardness is approximately 90 to 98 HRB If the welded assembly is given the customary double-aging heat treatment, the three areas identified above, as well as the unaffected base metal, transform and precipitation harden uniformly to a hardness range commensurate with the precipitation-hardening temperature. The weld metal may be somewhat less tough than the wrought base metal, as measured by the results of tensile-elongation and bend tests, depending on the type of joint, the welding process, and the hardening temperature. Higher precipitation-hardening temperature ensure good weld toughness with little sacrifice of strength. Maximum toughness requires the annealing of the weldment prior to the transformation and hardening treatments. Although other variations in the welding and heat-treating sequence are possible and may be desirable at times, the choice of the sequence should ensure that, after welding, the weld metal and the HAZ are in the annealed (austenitic) condition. To harden these areas, both the transformation and the precipitation-hardening heat treatments must be applied. If the components are given the transformation treatment before welding, the precipitation-hardening treatment alone, after welding, produces no significant hardening in either the weld metal or the HAZ. Austenitic Precipitation-Hardening Steels. The alloy content of these steels is high enough to maintain an austenitic structure after annealing and after any aging or hardening treatment. The precipitation-hardening phase is soluble at the annealing temperature of 1095 to 1120 °C (2000 to 2050 °F), and it remains in solution during rapid cooling from the annealing temperature. When these steels are reheated to about 650 to 760 °C (1200 to 1400 °F), precipitation occurs, and the hardness and strength of the austenitic structure increase. The hardness attained is lower than that of the martensitic or semiaustenitic PH steels, but the non-magnetic properties are retained. Although the austenitic PH steels remain austenitic during all phases of forming, welding, and heat treatment, some contain alloying elements (for precipitation-hardening purposes) that greatly affect behavior in welding. The austenitic precipitation-hardening stainless steels can be welded using the arc welding techniques described earlier in the section "Weldability" for the austenitic stainless steels. The major difference is that these steels are usually heat treated after welding to achieve the required mechanical properties, which is usually unnecessary with austenitic PH stainless steels. Austenitic precipitation-hardening stainless steels may be welded with matched or dissimilar filler metals or without filler metals, as is the case with most stainless steels. There is a wide variety of hardenable filler metals available for these PH grades through the manufacturer of consumables. The most commonly used grade is the 630 alloy, the only one currently included in the AWS specifications. Its composition is shown in Table 38 . Table 38 Chemical compositions of welding consumables for precipitation-hardening stainless steels All are maximum percentages. Composition, % AWS classification

C

Cr

Ni

Mo

Nb + Ta

Mn

Si

E630(a), ER630(b) 0.05 0.75 0.75 16.0−16.75 4.5−5.0 0.15−0.30 0.25−0.75 (a) Undiluted weld metal composition. (b) Consumable composition. Source: AWS A5.4−81, A5.9−81.

P

S

Cu

0.04

0.03

3.25−4.00

Duplex Stainless Steels (Ref 90, 91, 92). The physical and mechanical properties of duplex stainless steels affect the welding process. Because of their better stress-corrosion cracking resistance and appreciably higher yield and tensile strengths, these steels are currently used as direct substitutes for austenitic stainless steels when service above about 260 to 315 °C (500 to 600 °F) is not required.


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Duplex stainless steels are magnetic and have yield strengths typically about double that of type 316L, with tensile strengths considerably higher than those of standard austenitic grades. Duplex stainless steels also generally have a higher thermal conductivity and a lower coefficient of thermal expansion than austenitic stainless steels. Because of these factors, duplex stainless steels generally exhibit less distortion during welding than do austenite stainless steels. The general welding characteristics of the duplex alloy steels are very similar to those of austenitic stainless steels. They can be welded by any of the conventional arc welding processes. The normal arc welding processes, shielded metal arc, gas tungsten arc, gas metal arc, plasma arc, and submerged arc welded can all be used. In addition, electron beam and laser welding are used, as well as resistance welding. Heat input should be low enough to minimize intergranular carbide precipitation. Surface cleanliness is a must when welding duplex stainless steels. It is necessary to eliminate any source of hydrogen in the welding operation. For the gas-shielded processes, particularly on pipe, argon-helium purge gas should be used. Two differences between the characteristics of duplex stainless steels and those of austenitic stainless steels can be discerned: • Because duplex materials are appreciably stiffer than austenitic grades, greater forces are required for tube rolling into tube sheets, and so on • Because duplex alloys are highly sensitive to 475 °C (885 °F) and σ phase embrittlement control over interpass temperature is imperative It is good practice to weld austenitic grades at a maximum interpass temperature of approximately 200 °C (390 °F), but, in cases such as multipass orbital GTAW of pipe, interpass temperatures well above this level can often be tolerated with no significant adverse effect. These situations should be regarded with caution for duplex alloys, in view of embrittlement following prolonged exposure to temperatures much above 300 °C (570 °F), and a maximum interpass temperature of 200 °C (390 °F) is suggested. Wrought duplex stainless steels typically contain elongated islands of austenite in a ferrite matrix, as seen in Fig. 43 . Generally, weld metal contain less austenite than parent metal of the same composition. As-welded properties may be reduced because of this change in structure. Welding parameters and/or consumables are chosen to optimize the mechanical properties and corrosion resistance of the weld. Useful properties have been obtained using duplex stainless consumables with increased nickel to maintain sufficient austenite in the weld (Fig. 44 ). Slightly slower cooling rates after welding also permit more time for austenite formation and can improve as-welded mechanical properties and corrosion resistance. Weldments may also be annealed to optimize properties, but this is not always practical. It is essential that through cleaning be done after welding. Fig. 43 Photomicrograph of a duplex stainless steel showing elongated austenite islands in the ferrite matrix. The mill-annealed 19.1 mm (0.752 in.) thick plate sample is a longitudinal section etched using 15 ml HCl in 100 ml ethyl alcohol. 200×

Fig. 44 Photomicrograph of a duplex stainless weld obtained using consumables with increased nickel to retain sufficient austenite in the weld. Typical weld metal microstructure for gas metal arc weld in 19.1 mm (0.752 in.) thick plate. The light-etching phase is austenite. The micrograph shows a cross section etched in Groesbeck's reagent. 200×


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Test weldments may be evaluated in the laboratory before vessel fabrication and service exposure. The welding techniques, filler metal, and plate thickness studied in these evaluations are typically the same as those employed in field fabrication. Corrosion evaluation of weldments should employ environments and stress levels based on service expectations. Filler metal selection is critical. A filler metal with matching composition may result in a higher ferrite content than that of the base metal. Gooch (Ref 91) and others have shown that filler metals of duplex stainless steel with higher nickel and/or nitrogen are preferred. Weld metal cracking in duplex steels has rarely been identified. Because extensive austenite-to-ferrite and ferrite-to-austenite transformations occur in the HAZ, along with grain growth, the welding procedure must be carefully controlled. The extent of the above transformations depends to a great extent on the composition and the precise weld thermal cycle experienced by the HAZ during welding. Because of their high chromium level, these steels are prone to σ phase and 475 °C (885 °F) embrittlement. Though the weld thermal cycle is too short for σ-phase formation of 475 °C (885 °F) embrittlement to occur, care must be exercised in welding heavy-section steels of this type. REFERENCES 1. 2. 3. 4.

Metal Statistics: 1988, American Metal Market, Fairchild Publications, 1988 1988 Annual Statistical Report, American Iron and Steel Institute, 1989 D.J. De Renzo, Ed., Corrosion Resistant Materials Handbook, Noyes Data Corporation, 1985 R.L. Tobler, R.P. Reed, and D.S. Burkhalter, "Temperature Dependence of Yielding in Austenitic Stainless Steels," National Bureau of Standards, U.S. Department of Commerce 5. D.C. Larbalestier and H.W. King, Austenitic Stainless Steels at Cryogenic Temperatures, 1-Structural Stability and Magnetic Properties, Cryogenics, Vol 13 (No. 3), March 1973, p 160−168 6. K.R. Hanby et al., "Handbook on Materials for Superconducting Machinery," MCIC-HB-04, Metals and Ceramics Information Center, Battelle Columbus Laboratories, Jan 1977 7. A.J. Nachtigall, Strain Cycling Fatigue Behavior of Ten Structural Metals Tested in Liquid Helium, Liquid Nitrogen, and Ambient Air, in Properties of Materials for Liquified Natural Gas Tankage, STP 579, American Society for Testing and Materials, 1975, p 378−396 8. W. Weleff, H.S. McQueen, and W.F. Emmons, Cryogenic Tensile Properties of Selected Aerospace Materials, in Advances in Cryogenic Engineering, Vol 10, K.D. Timmerhaus, Ed., Plenum Press, 1965, p 14−15 9. L.P. Rice, J.E. Campbell, and W.F. Simmons, Tensile Behavior of Parent-Metal and Welded 5000-Series Aluminum Alloy Plate at Room and Cryogenic Temperatures, in Advances in Cryogenic Engineering, Vol 7, K.D. Timmerhaus, Ed., Plenum Press, 1962, p 478−489 10. K.A. Warren and R.P. Reed, Tensile and Impact Properties of Selected Materials from 20° to 300 °K, Monograph 63, National Bureau of Standards, U.S. Department of Commerce, June 1963 11. C.J. Guntner and R.P. Reed, Mechanical Properties of Four Austenitic Steels at Temperatures Between 300° and 20 °K, in Advances in Cryogenic Engineering, Vol 6, K.D. Timmerhaus, Ed., Plenum Press, 1961, p 565−576 12. C.J. Guntner and R.P. Reed, The Effect of Experimental Variables Including the Martensitic Transformation on the Low-Temperature Mechanical Properties of Austenitic Stainless Steels, Trans. ASM, Vol 55, Sept 1962, p 399−419 13. J.F. Watson and J.L. Christian, Low Temperature Properties of Cold-Rolled AISI Types 301, 302, 304ELC, and 310 Stainless Steel Sheet, in Low-Temperature Properties of High-Strength Aircraft and Missile Materials, STP 287, American Society for Testing and Materials, 1961, p 170−193 14. J.H. Bolton, L.L. Godby, and B.L. Taft, Materials for Use at Liquid Hydrogen Temperature, in Low-Temperature Properties of High-Strength Aircraft and Missile Materials, STP 287, American Society for Testing and Materials, 1961, p 108−120 15. H.L. Martin et al., "Effects of Low Temperature on the Mechanical Properties of Structural Metals," NASA SP-5012 (01),


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Office of Technology Utilization, National Aeronautics and Space Administration, 1968 16. D.T. Read and H.M. Ledbetter, Temperature Dependencies of the Elastic Constants of Precipitation-Hardened Aluminum Alloys 2014 and 2219, in J. Eng. Mater. Technol. (Trans. ASME), Series H, Vol 99 (No. 2), April 1977, p 181−184 17. J.M. Wells, W.A. Logsdon, and R. Kossowsky, Evaluations of Weldments in Austenitic Stainless Steels for Cryogenic Applications, in Advances in Cryogenic Engineering, Vol 24, K.D. Timmerhaus et al., Ed., Plenum Press, 1978, p 150−159 18. R.R. Vandervoort, Mechanical Properties of Inconel 625 Welds in 21-6-9 Stainless Steel, Cryogenics, Vol 18 (No. 8), Aug 1979, p 448−452 19. J.W. Montano, "The Stress Corrosion Resistance and the Cryogenic Temperature Mechanical Properties of Annealed Nitronic 60 Bar Material," NASA TM X-73359, National Aeronautics and Space Administration, Jan 1977 20. J.E. Campbell and L.P. Rice, Properties of Some Precipitation-Hardening Stainless Steels at Very Low Temperatures, in Low-Temperature Properties of High-Strength Aircraft and Missile Materials, STP 287, American Society for Testing and Materials, 1961, p 158−167 21. R.W. Finger, "Proof Test Criteria for Thin-Walled 2219 Aluminum Pressure Vessels," Vol I, NASA CR-135036, Vol II, NASA CR-135037, The Boeing Aerospace Company, Aug 1976 22. J.F. Watson and J.L. Christian, Mechanical Properties of High-Strength 301 Stainless Steel Sheet at 70, −320, and −423 F in Base Metal and Welded Joint Configuration, in Low-Temperature Properties of High-Strength Aircraft and Missile Materials, STP 287, American Society for Testing and Materials, 1961 23. H.M. Ledbetter, W.F. Weston, and E.R. Naimon, Low-Temperature Elastic Properties of Four Austenitic Stainless Steels, J. Appl. Phys., Vol 6 (No. 9), Sept 1975, p 3855−3860 24. R.L. Tobler et al., Low Temperature Fracture Behavior of Iron Nickel Alloy Steels, in Properties of Materials for Liquified Natural Gas Tankage, STP 579, American Society for Testing and Materials, Sept 1975, p 261−287 25. W.A. Logsdon, J.M. Wells, and R. Kossowsky, Fracture Mechanics Properties of Austenitic Stainless Steels for Advanced Applications, in Proceedings of the Second International Conference on Mechanical Behavior of Materials, American Society for Metals, 1976, p 1283−1289 26. R.P. Reed, R.L. Tobler, and R.P. Mikesell, The Fracture Toughness and Fatigue Crack Growth Rate of an Fe-Ni-Cr Superalloy at 298, 76, and 4K, in Advances in Cryogenic Engineering, Vol 22, K.D. Timmerhaus et al., Ed., Plenum Press, 1977, p 68−79 27. R.L. Tobler and R.P. Reed, Fatigue Crack Growth Resistance of Structural Alloys at Cryogenic Temperatures, in Advances in Cryogenic Engineering, Vol 24, K.D. Timmerhaus, et al., Ed., Plenum Press, 1978, p 82−90 28. T.F. Kiefer, R.D. Keys, and F.R. Schwartzberg, "Determination of Low-Temperature Fatigue Properties of Structural Metal Alloys," Final Report, The Martin Company, Oct 1965 29. E.H. Schmidt, "Fatigue Properties of Sheet, Bar, and Cast Metallic Materials for Cryogenic Applications," Report R-7564, Rocketdyne Division, North American Rockwell Corporation, Aug 1968 30. D.N. Gideon et al., The Fatigue Behavior of Certain Alloys in the Temperature Range from Room Temperature to −423 F, in Advances in Cryogenic Engineering, Vol 7, K.D. Timmerhaus, Ed., Plenum Press, 1962, p 503−508 31. D.N. Gideon et al., "Investigation of Notch Fatigue Behavior of Certain Alloys in the Temperature Range of Room Temperature to −423 F," ASD-TR-62-351, Battelle Memorial Institute, Aug 1962 32. B.F. Brown, Stress Corrosion Cracking Control Measures, Monograph 156, National Bureau of Standards, U.S. Department of Commerce, June 1977 33. R.M. Latanision and R.W. Staehle, Stress Corrosion Cracking of Iron-Nickel-Chromium Alloys, in Proceedings of Conference on Fundamental Aspects of Stress Corrosion Cracking, National Association of Corrosion Engineers, 1969, p 214−307 34. S.W. Dean, Review of Recent Studies on the Mechanism of Stress Corrosion Cracking in Austenitic Stainless Steels, in Stress Corrosion⎯New Approaches, STP 610, H.L. Craig, Jr., Ed., American Society for Testing and Materials, 1976, p 308−337 35. G. Fontana, Corrosion Engineering, 3rd ed., McGraw-Hill, 1986 36. Corrosion Resistance of the Austenitic Chromium-Nickel Stainless Steels in Atmospheric Environments, The International Nickel Company, Inc., 1963 37. K.L. Money and W.W. Kirk, Stress Corrosion Cracking Behavior of Wrought Fe-Cr-Ni Alloys in Marine Atmosphere, Mater. Perform., Vol 17, July 1978, p 28−36 38. M. Henthorne, T.A. DeBold, and R.J. Yinger, "Custom 450⎯A New High Strength Stainless Steel," Paper 53, presented at Corrosion/72, National Association of Corrosion Engineers, 1972 39. The Role of Stainless Steels in Desalination, American Iron and Steel Institute, 1974 40. M.A. Streicher, Analysis of Crevice Corrosion Data from Two Sea Water Exposure Tests on Stainless Alloys, Mater. Perform., Vol 22, May 1983, p 37−50 41. A.H. Tuthill and C.M. Schillmoller, Guidelines for Selection of Marine Materials, The International Nickel Company, Inc.,


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1971 42. R.M. Kain, "Crevice Corrosion Resistance of Austenitic Stainless Steels in Ambient and Elevated Temperature Seawater," Paper 230, presented at Corrosion/79, National Association of Corrosion Engineers, 1979 43. F.L. LaQue and H.R. Copson, Ed., Corrosion Resistance of Metals and Alloys, Reinhold, 1963, p 375−445 44. J.E. Truman, in Corrosion: Metal/Environmental Reactions, Vol 1, L.L. Shreir, Ed., Newness-Butterworths, 1976, p 352 45. M.A. Streicher, Development of Pitting Resistant Fe-Cr-Mo Alloys, Corrosion, Vol 30, 1974, p 77−91 46. H.O. Teeple, Corrosion by Some Organic Acids and Related Compounds, Corrosion, Vol 8, Jan 1952, p 14−28 47. T.A. DeBold, J.W. Martin, and J.C. Tverberg, Duplex Stainless Offers Strength and Corrosion Resistance, in Duplex Stainless Steels, R.A. Lula, Ed., American Society for Metals, 1983, p 169−189 48. L.A. Morris, in Handbook of Stainless Steels, D. Peckner and I.M. Bernstein, Ed., McGraw-Hill, 1977, p 17-1 49. "Material Requirements: Sulfide Stress Cracking Resistant Metallic Materials for Oil Field Equipment," MR-01-84, National Association of Corrosion Engineers 50. J.R. Kearns, M.J. Johnson, and J.F. Grubb, "Accelerated Corrosion in Dissimilar Metal Crevices," Paper 228, presented at Corrosion/86, National Association of Corrosion Engineers, 1986 51. L.S. Redmerski, J.J. Eckenrod, and K.E. Pinnow, "Cathodic Protection of Seawater-Cooled Power Plant Condensers Operating with High Performance Ferritic Stainless Steel Tubing," Paper 208, presented at Corrosion/85, National Association of Corrosion Engineers, 1985 52. E.C. Hoxie and G.E. Tuffnell, A Summary of INCO Corrosion Tests in Power Plant Flue Gas Scrubbing Processes, in Resolving Corrosion Problems in Air Pollution Control Equipment, National Association of Corrosion Engineers, 1976 53. Effective Use of Stainless Steel in FGD Scrubber Systems, American Iron and Steel Institute, 1978 54. G.T. Paul and R.W. Ross, Jr., "Corrosion Performance in FGD Systems at Laramie River and Dallman Stations," Paper 194, presented at Corrosion/83, National Association of Corrosion Engineers, 1983 55. A.P. Majidi and M.A. Streicher, "Four Non-Destructive Electrochemical Tests for Detecting Sensitization in Type 304 and 304L Stainless Steels," Paper 62, presented at Corrosion/85, National Association of Corrosion Engineers, 1985 56. A.H. Tuthill, Resistance of Highly Alloyed Materials and Titanium to Localized Corrosion in Bleach Plant Environments, Mater Perform., Vol 24, Sept 1985, p 43−49 57. "Standard Test Methods for Pitting and Crevice Corrosion Resistance of Stainless Steels and Related Alloys by the Use of Ferric Chloride Solution," G 48, Annual Book of ASTM Standards, American Society for Testing and Materials 58. "Standard Recommended Practice for Examination and Evaluation of Pitting Corrosion," G 46, Annual Book of ASTM Standards, American Society for Testing and Materials 59. T.A. DeBold, Which Corrosion Test for Stainless Steels, Mater. Eng., Vol 2 (No. 1), July 1980 60. R.M. Davison et al., A Review of Worldwide Developments in Stainless Steels in Specialty Steels and Hard Materials, Pergamon Press, 1983, p 67−85 61. A.M. Sabroff, F.W. Boulger, and H.J. Henning, Forging Materials and Practices, Reinhold, 1968 62. H.J. Henning, A.M. Sabroff, and F.W. Boulger, A Study of Forging Variables, Report ML-TDR-64-95, U.S. Air Force, 1964 63. Open Die Forging Manual, 3rd ed., Forging Industry Association, 1982, p 106−107 64. ASME Boiler and Pressure Vessel Code, Section III, Division I, Figure NB-2433.1-1, American Society of Mechanical Engineers, 1986 65. Machining and Abrasive Wheel Grinding of Carpenter Stainless Steels, in Carpenter Stainless Steels, Selection, Alloy Data, Fabrication, Carpenter Technology Corporation, 1987, p 240−241 66. V.A. Tipnis, Machining of Stainless Steels, Wire, Aug 1971, p 153−161 67. Metallurgy of Welding Stainless Steels, in Stainless Steels,,American Society for Metals, 1978, p 11-1 to 11−22 68. R.A. Lula, Fabrication of Stainless Steels⎯Machining, in Stainless Steels, American Society for Metals, 1986, p 112−114 69. "Free-Machining Stainless Steels," American Iron and Steel Institute, 1975 70. Guide to Machining Stainless Steels and Other Specialty Metals, Carpenter Technology Corporation, 1985 71. D.M. Blott, Machining Wrought and Cast Stainless Steels, in Handbook of Stainless Steels, McGraw-Hill, 1977, p 24-2 to 24−30 72. C.A. Divine, Jr., What to Consider in Choosing an Alloy, Met. Prog., Feb 1968, p 19−23 73. L. Colombier and J. Hochmann, Manufacturing, Forming and Finishing Techniques⎯Machining, in Stainless and Heat Resisting Steels, St. Martin's Press, 1968, p 508−514 74. Machining Operations, in Stainless Steel Fabrication, Allegheny Ludlum Steel Corporation, 1959, p 223−259 75. W.C. Clarke, Which Free-Machining Chromium Stainless?, Metalwork. Prod., Sept 1964, p 68−71 76. A. Moskowitz et al., Free-Machining Stainless Steels, U.S. Patent 3,401,035, 1968 77. F.M. Richmond, A Decade of Progress in Machinability, Finishing and Forming, Met. Prog., Aug 1967, p 85−86


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78. C.W. Kovach and A. Moskowitz, How to Upgrade Free-Machining Properties, Met. Prog., Aug 1967, p 173−180 79. J.R. Blank et al., Improved and More Consistent Steels for Machining, in Influence of Metallurgy on Machinability of Steel, Proceedings of an International Symposium, American Society for Metals, 1977, p 397−419 80. W.C. Clarke, Which Free-Machining Stainless?, Metalwork. Prod., 27 May 1964, p 43−45 81. J.A. Ferree, Jr., Free Machining Austenitic Stainless Steel, U.S. Patent 3,888,659, 1975 82. W.C. Clarke, Jr., Free-Machining Stainless Steel and Method, U.S. Patent 2,697,035, 1954 83. J.J. Eckenrod and C.W. Kovach, Effects of Manganese on Austenitic Stainless Steels, Met. Eng. Q., Feb 1972, p 5−10 84. R.P. Ney, Sr., Free Machining, Cold Formable Austenitic Stainless Steel, U.S. Patent 4,444,588, 1984 85. J.J. Eckenrod et al., Low Carbon Plus Nitrogen, Free-Machining Austenitic Stainless Steel, U.S. Patent 4,613,367, 1986 86. P.K. Wright and A. Bagchi, Wear Mechanisms That Dominate Tool-Life in Machining, J. Appl. Metalwork., Vol 1 (No. 4), 1981, p 15−23 87. "Carpenter Custom 630 (17Cr-4Ni)," Carpenter Technology Corporation, 1971 88. D. Pechner and I.M. Bernstein, Handbook of Stainless Steels, McGraw-Hill, 1977 89. Welding Handbook, Vol 4, Metals and Their Weldability, 9th ed., American Welding Society, 1982 90. H.B. Cary, Modern Welding Technology, 2nd ed., Prentice-Hall, 1989 91. T.G. Gooch, Weldability of Duplex Ferritic-Austenitic Stainless Steels, in Duplex Stainless Steels, R.A. Lula, Ed., Proceedings of Conference on Duplex Austenitic-Ferritic Stainless Steel, American Society for Metals, Oct 1982 92. S.A. David, Welding of Stainless Steels, in Encyclopedia of Materials Science and Engineering, Vol 7, M.B. Bever, Ed., The MIT Press, p 5316−5320 SELECTED REFERENCES • R.Q. Barr, Ed., Stainless Steel '77,Climax Molybdenum Company, 1978 • Book on Metals and Alloys in the Unified Numbering System, Society of Automotive Engineers Inc., and American Society for Testing and Materials • A.B. Kinzel et al., The Alloys of Iron and Chromium, 2 vol., McGraw-Hill, 1937 • R.A. Lula, Ed., New Developments in Stainless Steel Technology, in Proceedings of International Conference on New Developments in Stainless Steel Technology, American Society for Metals, 1985 • H.E. McGannon, Ed., The Making, Shaping and Treating of Steel, 9th ed., United States Steel Corporation, 1971 • J.H.G. Monypenny, Stainless Iron and Steel, 2 vol., Chapman and Hall, 1951 • J.G. Parr and A. Hanson, An Introduction to Stainless Steel, American Society for Metals, 1965 • D. Peckner and I.M. Bernstein, Handbook of Stainless Steels, McGraw-Hill, 1977 • F.B. Pickering, Ed., The Metallurgical Evolution of Stainless Steels, American Society for Metals, 1979 • A.O. Schaefer, Ed., Elevated Temperature Properties in Austenitic Stainless Steels, American Society of Mechanical Engineers, 1974 • A.J. Sedriks, Corrosion of Stainless Steels, Wiley-Interscience, 1979 • Source Book on Stainless Steels, American Society for Metals, 1976 • Stainless Steel for Architectural Use, STP 454, American Society for Testing and Materials, 1969 • Stainless Steels '87, The Institute of Metals, 1988 • E.E. Thum, Ed., The Book of Stainless Steels, American Society for Metals, 1935 • C.A. Zapffe, Stainless Steels, American Society for Metals, 1949


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Cast Stainless Steels

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Specialty Steels and Heat-Resistant Alloys Cast Stainless Steels Revised by Malcolm Blair, Steel Founders' Society of America STAINLESS STEELS are a class of chromium-containing steels widely used for their corrosion resistance in aqueous environments and for service at elevated temperatures. Stainless steels are distinguished from other steels by the enhanced corrosion and oxidation resistance created by chromium additions. Chromium imparts passivity to ferrous alloys when present in amounts of more than about 11%, particularly if conditions are strongly oxidizing. Consequently, steels with more than 10 or 12% Cr are sometimes defined as stainless steels. This article reviews the properties of cast steels that are specified either for liquid corrosion service at temperatures below 650 °C (1200 °F) or for service at temperatures above 650 °C (1200 °F). The cast steels suitable for these applications are often high-alloy compositions because carbon and low-alloy steels do not provide sufficient corrosion resistance and/or strength at elevated temperatures. These high-alloy cast steels generally have more than 10% Cr and primarily consist of stainless steel compositions. Stainless steel castings are usually classified as either corrosion-resistant castings (which are used in aqueous environments below 650 °C, or 1200 °F) or heat-resistant castings (which are suitable for service temperatures above 650 °C, or 1200 °F). However, this line of demarcation in terms of application is not always distinct, particularly for steel castings used in the range from 480 to 650 °C (900 to 1200 °F). The usual distinction between heat-resistant and corrosion-resistant cast steels is based on carbon content. In general, the cast and wrought stainless steels possess equivalent resistance to corrosive media and they are frequently used in conjunction with each other. Important differences do exist, however, between some cast stainless steels and their wrought counterparts. One significant difference is in the microstructure of cast austenitic stainless steels. There is usually a small amount of ferrite present in austenitic stainless steel castings, in contrast to the single-phase austenitic structure of the wrought alloys. The presence of ferrite in the castings is desirable for facilitating weld repair, but ferrite also increases resistance to stress-corrosion cracking. There have been only a few stress-corrosion cracking failures with cast stainless steels in comparison to the approximately equivalent wrought compositions. The principal reasons for this resistance are apparently: • Silicon added for fluidity gives added benefit from the standpoint of stress-corrosion cracking • Sand castings are usually tumbled or sandblasted to remove molding sand and scale; this probably tends to put the surface in compression Wrought and cast stainless steels may also differ in mechanical properties, magnetic properties, and chemical content. Because of the possible existence of large dendritic grains, intergranular phases, and alloy segregation, typical mechanical properties of cast stainless steels may vary more and generally are inferior to those of any wrought structure.

Grade Designations and Compositions 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 denotes the nominal chromium-nickel type of the alloy (Fig. 1 ). As nickel content increases, the second letter of the designation is changed from A to Z. The numeral or numerals following the first two letters indicate maximum carbon content (percentage × 100) of the alloy. 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 (Fig. 1 ), with a maximum carbon content of 0.08% and containing molybdenum (M). Fig. 1 Chromium and nickel contents in ACI standard grades of heat- and corrosion-resistant steel castings. See text for details.


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Some of the high-alloy cast steels exhibit many of the same properties of cast carbon and low-alloy steels (see the article"Steel Castings" in this Volume). Some of the mechanical properties of these grades (for example, hardness and tensile strength) can be altered by a suitable heat treatment. The cast high-alloy grades that contain more than 20 to 30% Cr plus Ni, however, do not show the phase changes observed in plain carbon and low-alloy steels during heating or cooling between room temperature and the melting point. These materials are therefore nonhardenable, and their properties depend on composition rather than heat treatment. Therefore, special consideration must be given to each grade of high-alloy cast steel with regard to casting design, foundry practice, and subsequent thermal processing (if any). Compositions of C-Type (Corrosion-Resistant) Steel Castings. The C-type steel castings for liquid corrosion service are often classified on the basis of composition, although it should be recognized that classification by composition often involves microstructural distinctions (see the section "Composition and Microstructure" in this article). Table 1 lists the compositions of the commercial cast corrosion-resistant alloys. Alloys are grouped as: • Chromium steels • Chromium-nickel steels, in which chromium is the predominant alloying element • Nickel-chromium steels, in which nickel is the predominant alloying element The serviceability of cast corrosion-resistant steels depends greatly on the absence of carbon, and especially precipitated carbides, in the alloy microstructure. Therefore, cast corrosion-resistant alloys are generally low in carbon (usually lower than 0.20% and sometimes lower than 0.03%). Table 1 Compositions and typical microstructures of ACI corrosion-resistant cast steels

ACI type

Composition, %(b)

Wroug ht alloy type(a)

ASTM specifications

Most common end-use microstructure

C

Mn

Si

Cr

Ni

Others(c)

Chromium steels CA-15

A 743, A 217, A 487

Martensite

0.15

1.00

1.50

11.5−14.0

1.0

0.50Mo(d)

...

A 743

Martensite

0.15

1.00

0.65

11.5−14.0

1.0

0.15−1.00Mo

CA-40

420

A 743

Martensite

0.40

1.00

1.50

11.5−14.0

1.0

0.5Mo(d)

CA-40F

...

A 743

Martensite

0.2−0.4

1.00

1.50

11.5−14.0

1.0

...

CB-30

431, 442

A 743

Ferrite and carbides

0.30

1.00

1.50

18.0−22.0

2.0

...

CC-50

446

A 743

Ferrite and carbides

0.30

1.00

1.50

26.0−30.0

4.0

...

CA-15M

410

Chromium-nickel steels CA-6N

...

A 743

Martensite

0.06

0.50

1.00

10.5−12.5 6.0−8.0

CA-6NM

...

A 743, A 487

Martensite

0.06

1.00

1.00

11.5−14.0 3.5−4.5 0.4−1.0Mo

CA-28MWV

...

A 743

Martensite

1.00

11.0−12.5 0.50−1. 0.9−1.25Mo; 00 0.9−1.5W; 0.2−0.3V

CB-7Cu-1

...

A 747

Martensite, age hardenable

0.07

0.70

1.00

15.5−17.7 3.6−4.6 2.5−3.2Cu; 0.20−0.35Nb; 0.05N max

CB-7Cu-2

...

A 747

Martensite, age hardenable

0.07

0.70

1.00

14.0−15.5 4.5−5.5 2.5−3.2Cu; 0.20−0.35Nb; 0.05N max

CD-4MCu

...

A 351, A 743, A 744, A 890

Austenite in ferrite, age hardenable

0.04

1.00

1.00

25.0−26.5 4.75−6. 1.75−2.25Mo; 0 2.75−3.25Cu


0.20−0.28 0.50−1. 00

...

Page 1410


CE-30

312

CF-3(e)


01 Sep 2005

A 743

Ferrite in austenite

0.30

1.50

2.00

26.0−30.0 8.0−11. 0

...

304L

A 351, A 743, A 744


0.03

1.50

2.00

17.0−21.0 8.0−12. 00

...

CF-3M(e)

316L

A 351, A 743, A 744


0.03

1.50

2.00

17.0−21.0 8.0−12. 2.0−3.0Mo 0

CF-3MN

...

A 743


0.03

1.50

1.50

17.0−21.0 9.0−13. 2.0−3.0Mo; 0 0.10−0.20N

CF-8(e)

304

A 351, A 743, A 744


0.08

1.50

2.00

18.0−21.0 8.0−11. 0

CF-8C

347

A 351, A 743, A 744


0.08

1.50

2.00

18.0−21.0 9.0−12. Nb(f) 0

CF-8M

316

A 351, A 743, A 744


0.8

1.50

2.00

18.0−21.0 9.0−12. 2.0−3.0Mo 0

CF-10

...

A 351


0.04−0.10

1.50

2.00

18.0−21.0 8.0−11. 0

CF-10M

...

A 351


0.04−0.10

1.50

1.50

18.0−21.0 9.0−12. 2.0−3.0Mo 0

CF-10MC

...

A 351


0.10

1.50

1.50

15.0−18.0 13.0−16 1.75−2.25Mo .0

CF-10SMnN

...

A 351, A 743


0.10

Ferrite in austenite or austenite

0.12

1.50

2.00

18.0−21.0 9.0−12. 2.0−3.0Mo 0

...

7.00−9. 3.50−4. 16.0−18.0 8.0−9.0 0.08−0.18N 00 50

CF-12M

316

CF-16F

303

A 743

Austenite

0.16

1.50

2.00

18.0−21.0 9.0−12. 1.50Mo max; 0 0.20−0.35Se

CF-20

302

A 743

Austenite

0.20

1.50

2.00

18.0−21.0 8.0−11. 0

A 351, A 743


0.06

4.00−6. 00

1.00

20.5−23.5 11.5−13 1.50−3.00Mo; .5 0.10−0.30Nb; 0.10−30V; 0.20−40N

A 351, A 743, A 744


0.08

1.50

1.50

18.0−21.0 9.0−13. 3.0−4.0Mo 0

CG-6MMN

...

...

...

...

CG-8M

317

CG-12

...

A 743


0.12

1.50

2.00

20.0−23.0 10.0−13 .0

...

CH-8

...

A 351


0.08

1.50

1.50

22.0−26.0 12.0−15 .0

...

CH-10

...

A 351


0.04−0.10

1.50

2.00

22.0−26.0 12.0−15 .0

...

CH-20

309

A 351, A 743

Austenite

0.20

1.50

2.00

22.0−26.0 12.0−15 .0

...

A 351, A 743, A 744


0.025

1.20

1.00

19.5−20.5 17.5−19 6.0−7.0V; .5 0.18−0.24N; 0.50−1.00Cu

A 743

Austenite

0.20

2.00

2.00

23.0−27.0 19.0−22 .0

CK-3MCuN

CK-20

...

310

...

Nickel-chromium steel CN-3M

...

A 743

Austenite

0.03

2.00

1.00

20.0−22.0 23.0−27 4.5−5.5Mo .0

CN-7M

...

A 351, A 743, A 744

Austenite

0.07

1.50

1.50

19.0−22.0 27.5−30 2.0−3.0Mo; 3.0−4.0 .5 Cu

CH-7MS

...

A 743, A 744

Austenite

0.07

1.50

CT-15C

...

A 351

Austenite


3.50(g) 18.0−20.0 22.0−25 2.5−3.0Mo; .0 1.5−2.0Cu

0.05−0.15 0.15−1. 0.50−1. 19.0−21.0 31.0−34 0.5−1.5V 50 50 .0

Page 1411



01 Sep 2005

(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); CF-10SMnN, 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); CA-28MWV, 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%.

All cast corrosion-resistant steels contain more than 11% chromium, and most contain from 1 to 30% nickel (a few have less than 1% Ni). About two-thirds of the corrosion-resistant steel castings produced in the United States are of grades that contain 18 to 22% Cr and 8 to 12% Ni. In general, the addition of nickel to iron-chromium alloys improves ductility and impact strength. An increase in nickel content increases resistance to corrosion by neutral chloride solutions and weakly oxidizing acids. The addition of molybdenum increases resistance to pitting attack by chloride solutions. It also extends the range of passivity in solutions of low oxidizing characteristics. The addition of copper to duplex (ferrite is austenite) nickel-chromium alloys produces alloys that can be precipitation hardened to higher strength and hardness. The addition of copper to single-phase austenitic alloys greatly improves their resistance to corrosion by sulfuric acid. In all iron-chromium-nickel stainless alloys, resistance to corrosion by environments that cause intergranular attack can be improved by lowering the carbon content. Information on the corrosion characteristics and mechanical properties of the C-type steel castings is provided in the section "Corrosion-Resistant Steel Castings" in this article. Compositions of H-Type (Heat-Resistant) Steel Castings. Castings are classified as heat resistant if they are capable of sustained operation while exposed, either continuously or intermittently, to operating temperatures that result in metal temperatures in excess of 650 °C (1200 °F). Heat-resistant steel castings resemble high-alloy corrosion-resistant steels except for their higher carbon content, which imparts greater strength at elevated temperature. The higher carbon content and, to a lesser extent, alloy composition ranges distinguish cast heat-resistant steel grades from their wrought counterparts. Table 2 summarizes the compositions of standard cast heat-resistant grades and three grade variations (HK30, HK40, HT30) specified in ASTM A 351 for elevated-temperature and corrosive service of pressure-containing parts. Table 2 Compositions of ACI heat-resistant casting alloys ACI designation HA

Composition, %(b)

UNS number ...

ASTM specifications(a) A 217

C

Cr

Ni

Si (max)

0.20 max

8−10

...

1.00

26−30

4 max

2.00

HC

J92605

A 297, A 608

0.50 max

HD

J93005

A 297, A 608

0.50 max

26−30

4−7

2.00

HE

J93403

A 297, A 608

0.20−0.50

26−30

8−11

2.00

HF

J92603

A 297, A 608

0.20−0.40

19−23

9−12

2.00

HH

J93503

A 297, A 608, A 447

0.20−0.50

24−28

11−14

2.00

HI

J94003

A 297, A 567, A 608

0.20−0.50

26−30

14−18

2.00

HK

J94224

A 297, A 351, A 567, A 608

0.20−0.60

24−28

18−22

2.00

HK30

...

A 351

0.25−0.35

23.0−27.0

19.0−22.0

1.75

HK40

...

A 351

0.35−0.45

23.0−27.0

19.0−22.0

1.75

HL

J94604

A 297, A 608

0.20−0.60

28−32

18−22

2.00

HN

J94213

A 297, A 608

0.20−0.50

19−23

23−27

2.00

HP

...

A 297

0.35−0.75

24−28

33−37

2.00

HP-50WZ(c)

...

0.45−0.55

24−28

33−37

2.50

HT

J94605

... A 297, A 351, A 567, A 608

0.35−0.75

13−17

33−37

2.50

HT30

...

A 351

0.25−0.35

13.0−17.0

33.0−37.0

2.50

HU

...

A 297, A 608

0.35−0.75

17−21

37−41

2.50

HW

...

A 297, A 608

0.35−0.75

10−14

58−62

2.50

HX ... A 297, A 608 2.50 035−0.75 15−19 64−68 (a) ASTM designations are the same as ACI designations. (b) Rem Fe 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).


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(c) Also contains 4 to 6% W, 0.1 to 1.0% Zr, and 0.035% S (max) and P (max)

The three principal categories of H-type cast steels, based on composition, are: • Iron-chromium alloys • Iron-chromium-nickel alloys • Iron-nickel-chromium alloys Information on the properties of H-type grades of steel castings is contained in the section "Heat-Resistant Cast Steels" in this article.

Composition and Microstructure As shown in Table 1 , cast stainless steels can also be classified on the basis of microstructure. Structures may be austenitic, ferritic, martensitic, or ferritic-austenitic (duplex). The structure of a particular grade is primarily determined by composition. Chromium, molybdenum, and silicon promote the formation of ferrite (magnetic), while carbon, nickel, nitrogen, and manganese favor the formation of austenite (nonmagnetic). For example, a cast extra-low-carbon grade such as 0.03% C (max) cannot be completely nonmagnetic unless it contains 12 to 15% Ni. The wrought grades of these alloys normally contain about 13% Ni. They are made fully austenitic to improve rolling and forging characteristics. Chromium (a ferrite and martensite promoter), nickel, and carbon (austenite promoters) are particularly important in determining microstructure (see the section "Ferrite Control " in this article). In general, straight chromium grades of high-alloy cast steel are either martensitic or ferritic, the chromium-nickel grades are either duplex or austenitic, and the nickel-chromium steels are fully austenitic. Ferrite in Cast Austenitic Stainless Steels. Cast austenitic alloys usually have from 5 to 20% ferrite distributed in discontinuous pools throughout the matrix, the percent of ferrite depending on the nickel, chromium, and carbon contents (see the section "Ferrite Control" ). The presence of ferrite in austenite may be beneficial or detrimental, depending on the application. Ferrite is beneficial and intentionally present in various corrosion-resistant cast steels (see some of the CF grades in Table 1 , for example) to improve weldability and to maximize corrosion resistance in specific environments. Ferrite is also used for strengthening duplex alloys. The section "Austenitic-Ferritic (Duplex) Alloys" in this article gives further information. Ferrite can be beneficial in terms of weldability because fully austenitic stainless steels are susceptible to a weldability problem known as hot cracking, or microfissuring. The intergranular cracking occurs in the weld deposit and/or in the weld heat-affected zone and can be avoided if the composition of the filler metal is controlled to produce about 4% ferrite in the austenitic weld deposit. Duplex CF grade alloy castings are immune to this problem. The presence of ferrite in duplex CF alloys improves the resistance to stress-corrosion cracking (SCC) and generally to intergranular attack. In the case of SCC, the presence of ferrite pools in the austenite matrix is thought to block or make more difficult the propagation of cracks. In the case of intergranular corrosion, ferrite is helpful in sensitized castings because it promotes the preferential precipitation of carbides in the ferrite phase rather than at the austenite grain boundaries, where they would increase susceptibility to intergranular attack. The presence of ferrite also places additional grain boundaries in the austenite matrix, and there is evidence that intergranular attack is arrested at austenite-ferrite boundaries. It is important to note, however, that not all studies have shown ferrite to be unconditionally beneficial to the general corrosion resistance of cast stainless steels. Some solutions attack the austenite phase in heat-treated alloys, whereas others attack the ferrite. For instance, calcium chloride solutions attack the austenite. On the other hand, a 10 ° Baumé cornstarch solution, acidified to a pH of 1.8 with sulfuric acid and heated to a temperature of 135 °C (275 °F), attacks the ferrite. Whether corrosion resistance is improved by ferrite and to what degree depends on the specific alloy composition, the heat treatment, and the service conditions (environment and stress state). Ferrite can be detrimental in some applications. One concern may be the reduced toughness from ferrite, although this is not a major concern, given the extremely high toughness of the austenite matrix. A much greater concern is for applications that require exposure to elevated temperatures, usually 315 °C (600 °F) and higher, where the metallurgical changes associated with the ferrite can be severe and detrimental. In applications requiring that these steels be heated in the range from 425 to 650 °C (800 to 1200 °F), carbide precipitation occurs at the edges of the ferrite pools in preference to the austenite grain boundaries. When the steel is heated above 540 °C (1000 °F), the ferrite pools transform to a χ or σ phase. If these pools are distributed in such a way that a continuous network is formed, embrittlement or a network of corrosion penetration may result. Also, if the amount of ferrite is too great, the ferrite may form continuous stringers where corrosion can take place, producing a condition similar to grain boundary attack. In the lower end of this temperature range, the reductions in toughness observed have been attributed to carbide precipitation or reactions associated with 475 °C (885 °F) embrittlement. The 475 °C embrittlement is caused by the precipitation of an intermetallic phase will a composition of approximately 80Cr-20Fe. The name derives from the fact that this embrittlement is most severe and rapid when it occurs at approximately 475 °C (885 °F). At 540 °C (1000 °F) and above, the ferrite phase may transform to a complex iron-chromium-nickel-molybdenum intermetallic compound known as σ phase, which reduces toughness, corrosion resistance, and creep ductility. The extent of the reduction


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increases with time and temperature to about 815 °C (1500 °F) and may persist to 925 °C (1700 °F). In extreme cases, Charpy V-notch energy at room temperature may be reduced 95% from its initial value (Ref 1, 2). At temperatures above 540 °C (1000 °F), austenite also has better creep resistance than ferrite. The weaker ferrite phase may lend better plasticity to the alloy, but after long exposure at temperatures in the 540 to 760 °C (1000 to 1400 °F) range, it may transform to σ or χ phase, which reduces resistance to impact. In some instances, the alloy is deliberately aged to form the σ or χ phase and thus increase strength. Austenite can transform directly to σ or χ without going through the ferrite phase. In weld deposits, the presence of σ or χ phase is extremely detrimental to ductility When welding for service at room temperature or up to 540 °C (1000 °F), 4 to 10% ferrite may be present and will greatly reduce the tendency toward weld cracking. However, for service at temperatures between 540 and 815 °C (1000 and 1500 °F), the amount of ferrite in the weld must be reduced to less than 5% to avoid embrittlement from excessive σ or χ phase. Ferrite Control. From the preceding discussion, it is apparent that ferrite in predominantly austenitic cast stainless steels can offer property advantages in some steels (notably the CF alloys) and disadvantages in other cases (primarily at elevated temperatures). The underlying causes for the dependence of ferrite content on composition are found in the phase equilibria for the iron-chromium-nickel system. These phase equilibria have been exhaustively documented and related to commercial stainless steels. The major elemental components of cast stainless steels are in competition to promote austenite or ferrite phases in the alloy microstructure. Chromium, silicon, molybdenum, and niobium promote the presence of ferrite in the alloy microstructure; nickel, carbon, nitrogen, and manganese promote the presence of austenite. By balancing the contents of ferrite-and austenite-forming elements within the specified ranges for the elements in a given alloy, it is possible to control the amount of ferrite present in the austenite matrix. The alloy can usually be made fully austenitic or with ferrite contents up to 30% or more in the austenite matrix. The relationship between composition and microstructure in cast stainless steels permits the foundryman to predict and control the ferrite content of an alloy, as well as its resultant properties, by adjusting the composition of the alloy. This is accomplished with the Schoefer constitution diagram for cast chromium-nickel alloys (Fig. 2 ). This diagram was derived from an earlier diagram developed by Schaeffler for stainless steel weld metal (Ref 1). The use of Fig. 2 requires that all ferrite-stabilizing elements in the composition be converted into chromium equivalents and that all austenite-stabilizing elements be converted into nickel equivalents by means of empirically derived coefficients representing the ferritizing or austenitizing power of each element. A composition ratio is then obtained from the total chromium equivalent, Cre, and nickel equivalent, Nie, calculated for the alloy composition by: Cre = %Cr + 1.5(%Si) + 1.4(%Mo) Ã+ %Nb − 4.99 (Eq 1) Nie = %Ni + 30(%C) + 0.5 (%Mn) Ã+ 26(%N − 0.02) + 2.77 (Eq 2) where the elemental concentrations are given in weight percent. Although similar expressions have been derived that take into account additional alloying elements and different compositional ranges in the iron-chromium-nickel alloy system, use of the Schoefer diagram has become standard for estimating and controlling ferrite content in stainless steel castings. Fig. 2 Schoefer diagram for estimating the ferrite content of steel castings in the composition range of 16 to 26% Cr, 6 to 14% Ni, 4% Mo (max), 1% Nb (max), 0.2% C (max), 0.19% N (max), 2% Mn (max), and 2% Si (max). Dashed lines denote scatter bands caused by the uncertainty of the chemical analysis of individual elements. See text for equations used to calculate Cre and Nie. Source: Ref 1


Page 1414



01 Sep 2005

The Schoefer diagram possesses obvious utility for casting users and the foundryman. It is helpful for estimating or predicting ferrite content if the alloy composition is known and for setting nominal values for individual elements when calculating the furnace charge for an alloy in which a specified ferrite range is desired. Limits of Ferrite Control. Although ferrite content can be estimated and controlled on the basis of alloy composition only, there are limits to the accuracy with which this can be done. The reasons for this are many. First, there is an unavoidable degree of uncertainly in the chemical analysis of an alloy (note the scatter band in Fig. 2 ). Second, in addition to composition, the ferrite content depends on thermal history, although to a lesser extent. Third, ferrite contents at different locations in individual castings can vary considerably, depending on section size, ferrite orientation, presence of alloying-element segregation, and other factors. Both the foundryman and the user of stainless steel castings should recognize that the factors mentioned above place significant limits on the degree to which ferrite content (either as ferrite number or ferrite percentage) can be specified and controlled. In general, the accuracy of ferrite measurement and the precision of ferrite control diminish as the ferrite number increases. As a working rule, it is suggested that the ±6 about the mean or desired ferrite number be viewed as a limit of ferrite control under ordinary circumstances, with ±3 possible under ideal circumstances.

Heat Treatment The heat treatment of stainless steel castings is very similar in purpose and procedure to the thermal processing of comparable wrought materials (see the article "Heat Treating of Stainless Steels" in Heat Treating, Volume 4 of ASM Handbook. However, some differences warrant separate consideration here. Homogenization. Alloy segregation and dendritic structures may occur in castings and may be particularly pronounced in heavy sections. Because castings are not subjected to the high-temperature mechanical reduction and soaking treatments involved in the mill processing of wrought alloys, it is frequently necessary to homogenize some alloys at temperatures above 1095 °C (2000 °F) to promote uniformity of chemical composition and microstructure. The full annealing of martensitic castings results in recrystallization and maximum softness, but it is less effective than homogenization in eliminating segregation. Homogenization is a common procedure in the heat treatment of precipitation-hardening castings. Sensitization and Solution Annealing of Austenitic and Duplex Alloys. When austenitic or duplex (ferrite in austenite matrix) stainless steels are heated in or cooled slowly through a temperature range of about 425 to 870 °C (800 to 1600 °F), chromium-rich carbides form at grain boundaries in austenitic alloys and at ferrite-austenite interfaces in duplex alloys. These carbides deplete the surrounding matrix of chromium, thus diminishing the corrosion resistance of the alloy. In small amounts, these carbides may lead to localized pitting in the alloy, but if the chromium-depleted zones are extensive throughout the alloy or heat-affected zone (HAZ) of a weld, the alloy may disintegrate intergranularly in some environments. An alloy in this condition of reduced corrosion resistance due to the formation of chromium carbides is said to be sensitized, a situation that is most pronounced for austenitic alloys. In austenitic structures, the complex chromium carbides precipitate preferentially along the grain boundaries. This microstructure is susceptible to intergranular corrosion, especially in oxidizing solutions. In partially ferritic alloys, carbides tend to precipitate in the discontinuous carbide pools; thus, these alloys are less susceptible to intergranular attack. Solution annealing of austenitic and duplex stainless steels makes these alloys less susceptible to intergranular attack by ensuring the complete solution of the carbides in the matrix. Depending on the specific alloy in question, temperatures between 1040 and 1205 °C (1900 and 2200 °F) will ensure the complete solution of all carbides and phases, such as σ and χ, that sometimes form in highly alloyed stainless steels. Alloys containing relatively high total alloy content, particularly high molybdenum content, often require the higher solution treatment temperature. Water quenching from the temperature range of 1040 to 1205 °C (1900 to 2200 °F) normally completes the solution treatment. Solution-annealing procedures for all austenitic


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alloys require holding for a sufficient amount of time to accomplish the complete solution of carbides and quenching at a rate fast enough to prevent reprecipitation of the carbides, particularly while cooling through the range of 870 to 540 °C (1600 to 1000 °F). A two-step heat-treating procedure can be applied to the niobium-containing CF-8C alloy. The first treatment consists of solution annealing. This is followed by a stabilizing treatment at 870 to 925 °C (1600 to 1700 °F), which precipitates niobium carbides, prevents the formation of damaging chromium carbides, and provides maximum resistance to intergranular attack. Because of their low carbon content, CF-3 and CF-3M as-cast do not contain enough chromium carbide to cause selective intergranular attack; therefore, these alloys can be used in some environments in this condition. However, for maximum corrosion resistance, these grades require solution annealing. If the usual quenching treatment is difficult or impossible, holding for 24 to 48 h at 870 to 980 °C (1600 to 1800 °F) and air cooling is helpful for improving the resistance of castings to intergranular corrosion. However, except for alloys of very low carbon content and castings with thin sections, this treatment fails to produce material with as good a resistance to intergranular corrosion as properly quench-annealed material.

Corrosion-Resistant Steel Castings As previously mentioned, various high-alloy steel castings are classified as corrosion resistant (Table 1 ). These corrosion-resistant cast steels are widely used in chemical processing and power-generating equipment that requires corrosion resistance in aqueous or liquid-vapor environments at temperatures normally below 315 °C (600 °F). These alloys are also used in special applications with temperatures up to 650 °C (1200 °F). Compositions The chemical compositions of various corrosion-resistant cast steels are given in Table 1 . These cast steels are specified in the ASTM standards listed in Table 1 . Straight chromium stainless steels contain 10 to 30% Cr and little or no nickel. Although about two-thirds of the corrosion-resistant steel castings produced in the United States are of grades that contain 18 to 22% Cr and 8 to 12% Ni, the straight chromium compositions are also produced in considerable quantity, particularly the steel with 11.5 to 14.0% Cr. Corrosion resistance improves as chromium content is increased. In general, intergranular corrosion is less of a concern in the straight chromium alloys (which are typically ferritic), especially those containing 25% Cr or more. This is attributed to the high bulk chromium contents and the rapid diffusion rates of chromium in ferrite. Iron-chromium-nickel alloys have found wide acceptance and constitute about 60% of total production of high-alloy castings. They generally are austenitic with some ferrite. The most popular alloys of this type are CF-8 and CF-8M. these alloys are nominally 18-8 stainless steels are the cast counterparts of wrought types 304 and 316, respectively. The carbon content of each is maintained at 0.08% (max). Effects of Molybdenum on Corrosion Resistance. Alloys CF-3M and CF-8M are modifications of CF-3 and CF-8 containing 2 to 3% Mo to enhance general corrosion resistance. Their passivity under weakly oxidizing conditions is more stable than that of CF-3 and CF-8. The addition of 2 to 3% Mo increases resistance to corrosion by seawater and improves resistance to many chloride-bearing environments. The presence of 2 to 3% Mo also improves crevice corrosion and pitting resistance compared to the CF-8 and CF-3 alloys. The CF-8M and CF-3M alloys have good resistance to such corrosive media as sulfurous and acetic acids and are more resistant to pitting by mild chlorides. These alloys are suitable for use in flowing seawater, but will pit under stagnant conditions. Alloy CG-8M is slightly more highly alloyed than the CF-8M alloys, with the primary addition being increased molybdenum (3 to 4%). The increased amount of molybdenum provides superior corrosion resistance to halide-bearing media and reducing acids, particularly H2SO3 and H2SO4 solutions. The high molybdenum content, however, renders CG-8M generally unsuitable in highly oxidizing environments. Molybdenum-bearing alloys are generally not as resistant to highly oxidizing environments (this is particularly true for boiling HNO3), but for weakly oxidizing environments and reducing environments, Mo-bearing alloys are generally superior. Molybdenum may also produce detrimental catalytic reactions. For example, the residual molybdenum in CF-8 alloy must be held below 0.5% in the presence of hydrazine. Effects of Chromium, Carbon, and Silicon on Corrosion Resistance. In alloys of the CF type, the effects of composition on rates of general corrosion attack have been studied, and certain definite relationships have been established. Through the use of the Huey test (five 48 h periods of exposure to boiling 65% nitric acid, as described in practice C of ASTM A 262), it has been shown that, in this standardized environment, carbide-free quench-annealed alloys of various nickel, chromium, silicon, carbon, and manganese contents have corrosion rates directly related to these contents. Figure 3 shows the influences on corrosion rate exerted by various elements in a 19Cr-9Ni casting alloy. Variations in nickel, manganese, and nitrogen contents for the ranges shown have relatively slight influences, but variations in chromium, carbon, and silicon have marked effects. The relationship between composition and corrosion rate for properly heat-treated CF alloys in boiling 65% nitric acid is summarized in the nomograph presented in Fig. 4 . Fig. 3 Effects of various elements in a 19Cr-9Ni casting alloy on corrosion rate in boiling 65% nitric acid. Data were determined for solution-annealed and quenched specimens. Composition of base alloy was 19Cr, 9Ni, 0.09C, 0.8Mn, 1.0Si,


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0.04P (max), 0.03S (max), 0.06N.

Fig. 4 Nomograph for determining corrosion rate in boiling 65% nitric acid for solution-annealed and quenched type CF casting alloys

Iron-Nickel-Chromium Alloys. For some types of service, extensive use is made of iron-nickel-chromium alloys that contain more nickel than chromium. Most important among this group is alloy CN-7M, which has a nominal composition of 28% Ni, 20% Cr, 3.5% Cu, 2.5% Mo, and 0.07% C (max). In effect, this alloy is made by adding 20% Ni and 3.5% Cu to alloy CF-8M, which greatly improves resistance to hot, concentrated, weakly oxidizing solutions such as sulfuric acid and also improves resistance to severely oxidizing media. Alloys of this type can withstand all concentrations of sulfuric acid at temperatures up to 65 °C (150 °F) and many concentrations up to 80 °C (175 °F). They are widely used in nitric-hydrofluoric pickling solutions; phosphoric acid; cold dilute hydrochloric acid; hot acetic acid; strong, hot caustic solutions; brines; and many complex plating solutions and rayon spin baths. Results of in-plant corrosion testing of CF-8, CF-8M, and CN-7M alloys are shown in Table 3 . These tests give the specific effect of molybdenum on 19Cr-9Ni alloys in reducing selective attack and pitting, and the overall corrosion rate computed from loss in weight. The higher nickel plus copper and molybdenum in the CN-7M alloy reduces the rate of corrosion to a rate lower than that of the CF-8M alloy. Table 3 Results of in-plant corrosion testing of CF-8, CF-8M, and CN-7M alloys Temperat ure of


Metal loss on

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Type and composition of corroding solution


solution, °C

°F

100

212 CF-8

Settling tank after neutralizer: ammonium sulfate plus excess of sulfuric acid

50

99 to 100% fuming nitric acid

Saturated solution of sodium chloride plus 15% sodium sulfate; pH of 4.5

50

20

60

Alloy

665

26.2

Very heavy etch(a)

CF-8M

28

1.1

Light tarnish(b)

CN-7M

18

0.7

Bright

385

15.2

Very heavy etch(a)

CF-8M

10

0.4

Slight tarnish

CN-7M

2.5

0.1

Bright(b)

122 CF-8

122 CF-8

68

Surface condition by visual µm/yr mils/yr examination surface

Neutralizer after formation of ammonium sulfate: ammonium sulfate plus small excess of sulfuric acid, ammonia vapor, and steam

Ammonium sulfate processing solution: ammonium sulfate at pH of 8.0

01 Sep 2005

685

27.0

Heavy etch

CF-8M

175

6.8

Moderate etch

CN-7M

50

2.0

Light etch

CF-8

245

9.6

Moderate etch

CN-7M

79

3.1

Light etch

CF-8M

345

13.5

Moderate etch

2.5

0.1

Bright

240

9.5

140 CF-8M

Remarks CF-8M was installed for low corrosion tolerance equipment in this service and performed satisfactorily CF-8 in service showed excessive corrosion rate plus heavy concentration cell attack CF-8M had too high a corrosion rate in service for good valve life, although suitable for equipment of greater corrosion tolerance. CN-7M was installed in this service CF-8 was satisfactory except for low-tolerance equipment such as valves. CN-7M valves performed satisfactorily in service

CF-8M was installed for valves in service Concentration cell corrosion at various small areas of specimen (a) Concentration cell attack under insulating washer. (b) Slight concentration cell attack under insulating washer CF-8

Corrosion From Chlorine. The influence of contaminants is one of the most important considerations in selecting an alloy for a particular process application. Ferric chloride in relatively small amounts, for example, will cause concentration cell corrosion and pitting. The buildup of corrosion products in a chloride solution may increase the iron concentration to a level high enough to be destructive. Thus, chlorine salts, wet chlorine gas, and unstable chlorinated organic compounds cannot be handled by any of the iron-base alloys, creating a need for nickel-base alloys. Microstructures Although corrosion-resistant cast steels are usually classified on the basis of composition, it should be recognized that classification by composition also often involves microstructural distinctions. Table 1 shows the typical microstructures of various corrosion-resistant cast steels. As noted previously, straight chromium grades of high-alloy cast steel are either martensitic or ferritic, the chromium-nickel grades are either duplex or austenitic, and the nickel-chromium steels are fully austenitic. Martensitic grades include Alloys CA-15, CA-40, CA-15M, and CA-6NM. The CA-15 alloy contains the minimum amount of chromium necessary to make it essentially rustproof. It has good resistance to atmospheric corrosion, as well as to many organic media in relatively mild service. A higher-carbon modification of CA-15, CA-40 can be heat treated to higher strength and hardness levels. Alloy CA-15M is a molybdenum-containing modification of CA-15 that provides improved elevated-temperature strength. Alloy CA-6NM is an iron-chromium-nickel-molybdenum alloy of low carbon content. Austenitic grades include CH-20, CK-20 and CN-7M. The CH-20 and CK-20 alloys are high-chromium, high carbon, wholly austenitic compositions in which the chromium content exceeds the nickel content. The more highly alloyed CN-7M, as described earlier in the section "Iron-Nickel-Chromium Alloys," has excellent corrosion resistance in many environments and is often used in sulfuric acid environments. The CN-7MS alloy has a corrosion resistance similar to that of CN-7M. The CN-7MS alloy has outstanding resistance to corrosion from high-strength (>90%) nitric acid. Ferritic grades include CB-30 and CC-50. Alloy CB-30 is practically nonhardenable by heat treatment. As this alloy is normally made, the balance among the elements in the composition results in a wholly ferritic structure similar to wrought AISI type 442 stainless steel. Alloy CC-50 has substantially more chromium than CB-30 and has relatively high resistance to localized corrosion in many environments. Austenitic-ferritic (duplex) alloys include CE-30, CF-3, CF-3A, CF-8, CF-8A, CF-20, CF-3M, CF-3MA, CF-8M, CF-8C, CF-16F, and CG-8M. The microstructures of these alloys usually contain 5 to 40% ferrite, depending on the particular grade and the balance among the ferrite-promoting and austenite-promoting elements in the chemical composition (see the section "Ferrite Control" in this article). Duplex alloys offer superior strength, corrosion resistance, and weldability. The use of duplex cast steels has focused primarily on the CF grades, particularly by the power generation industry. Strengthening in the cast CF grade alloys is limited essentially to that which can be gained by incorporating ferrite into the austenite matrix phase. These alloys cannot be strengthened by thermal treatment, as can the cast martensitic alloys, not by hot or


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01 Sep 2005

cold working, as can the wrought austenitic alloys. Strengthening by carbide precipitation is also out of the question because of the detrimental effect of carbides on corrosion resistance in most aqueous environments. Thus, the alloys are effectively strengthened by balancing the alloy composition to produce a duplex microstructure consisting of ferrite (up to 40% by volume) distributed in an austenite matrix. It has been shown that the incorporation of ferrite into 19Cr-9NI cast steels improves yield and tensile strengths without substantial loss of ductility or impact toughness at temperatures below 425 °C (800 °F). The magnitude of this strengthening effect for CF-8 and CF-8M alloys at room temperature is shown in Fig. 5 . Table 4 shows the effect of ferrite content on the tensile properties of 19Cr-9Ni alloys at room temperature and at 355 °C (670 °F). Table 5 shows the effect of ferrite content on impact toughness. Table 4 Effect of ferrite content on tensile properties of 19Cr-9Ni alloys Tensile strength

Yield strength at 0.2% offset

MPa

ksi

MPa

ksi


Reduction in area, %

3

465

67.4

216

31.3

60.5

64.2

10

498

72.2

234

34.0

61.0

73.0

20

584

84.7

296

43.0

53.5

58.5

41

634

91.9

331

48.0

45.5

47.9

3

339

49.1

104

15.1

45.5

63.2

10

350

50.8

109

15.8

43.0

69.7

20

457

66.3

183

26.5

36.5

47.5

41

488

70.8

188

27.3

33.8

49.4

Ferrite content, % Tested at room temperature

Tested at 355 °C (670 °F)

Table 5 Charpy V-notch impact energy, ferrite content, and Cre/Nie ratio of duplex cast steels Charpy V-notch energy

Ferrite content, % MG(a)

FS(b)

Cre/Nie ratio(c)

28.5

20

20

1.5

20.7

12.5

14

1.4

Alloy

J

ft · lbf

Calculated

CF 3M

197

145

CF 3C

183

135

CG 8M

216

159

18

9

10

1.34

CF 3C

>358

>264

15

13

15

1.29

7

1.12

CF 3M >358 >264 7.7 6 (a) MG, magna gage. (b) FS, ferrite scope. (c) See Eq 1 and 2 for formulas to compute Cre and Nie.

Fig. 5 Yield strength and tensile strength versus percentage of ferrite for CF-8 and CF-8M alloys. Curves are mean values for 277 heats of CF-8 and 62 heats of CF-8M. Source: Ref 3

Other duplex alloys of interest include CD-4MCu and Ferralium. Alloy CD-4MCu is the most highly alloyed duplex alloy. Ferralium was developed by Langley Alloys and is essentially CD-4MCu with about 0.15% N added. With high levels of ferrite


Page 1419



01 Sep 2005

(about 40 to 50%) and low nickel, the duplex alloys have better resistance to stress-corrosion cracking (SCC) than CF-3M. Alloy CD-4MCu, which contains no nitrogen and has a relatively low molybdenum content, has only slightly better resistance to localized corrosion than CF-3M. Ferralium which has nitrogen and slightly higher molybdenum than CD-4MCu, exhibits better localized corrosion resistance than either CF-3M or CD-4MCu. Improvements in stainless steel production practices (for example, electron beam refining, vacuum and argon-oxygen decarburization, and vacuum induction melting) have also created a second generation of duplex stainless steels. These steels offer excellent resistance to pitting and crevice corrosion, significantly better resistance to chloride SCC than the austenitic stainless steels, good toughness, and yield strengths two to three higher than those of type 304 or 316 stainless steels. First-generation duplex stainless steels, for example, AISI type 329 and CD-4MCu, have been in use for many years. The need for improvement in the weldability and corrosion resistance of these alloys resulted in the second-generation alloys, which are characterized by the addition of nitrogen as an alloying element. Second-generation duplex stainless steels are usually about a fifty-fifty blend of ferrite and austenite. The new duplex alloys combine the near immunity to chloride SCC of the ferritic grades with the toughness and ease of fabrication of the austenitics. Among the second-generation duplexes, Alloy 2205 seems to have become the general-purpose stainless. Table 6 lists the nominal compositions of first- and second-generation duplex alloys. Table 6 Nominal compositions of first- and second-generation duplex stainless steels Composition, %(a) UNS designation

Common name

Cr

Ni

Mo

Cu

N

Others

18.5

4.7

2.7

...

...

1.7Si

First generation steels S31500

3RE60

S32404

Uranus 50

21

7.0

2.5

1.5

...

...

S32900

Type 329

26

4.5

1.5

...

...

...

J93370

CD-4MCu

25

5

2

3

...

...

Second generation steels S31200

44LN

25

6

1.7

...

0.15

...

S31260

DP-3

25

7

3

0.5

0.15

0.3W

S31803

Alloy 2205

22

5

3

...

0.15

...

S32550

Ferralium 255

25

6

3

2

0.20

...

S32950

7-Mo PLUS

26.5

4.8

1.5

...

0.20

...

25

7

4.5

...

0.25

...

J93404 Atlas 958, COR 25 (a) All compositions contain balance of iron.

Precipitation-Hardening Alloys. Corrosion-resistant alloys capable of being hardened by low-temperature treatment to obtain improved mechanical properties are usually duplex-structure alloys with much more chromium than nickel. The addition of copper enables these alloys to be strengthened by precipitation hardening. These alloys are significantly higher in strength than the other corrosion-resistant alloys even without hardening. The alloys CB-7Cu-1 and CB-7Cu-2 have corrosion resistances between those of CA-15 and CF-18. They are widely used for structural components requiring moderate corrosion resistance, as well as for components requiring resistance to erosion and wear. The alloy CD-4MCu is widely used in many applications where its good corrosion resistance (which often equals or even exceeds that of CF-8M) and excellent resistance to erosion make it the most desirable alloy. The steel CD-4MCu has outstanding resistance to nitric acid and mixtures of nitric acid and organic acids, as well as excellent resistance to a wide range of corrosive chemical process conditions. This alloy is normally used in the solution-annealed condition, but it can be precipitation hardened for carefully selected applications when lower corrosion resistance can be tolerated and when there is no potential for stress-corrosion cracking. Corrosion Characteristics Table 7 compares the general corrosion resistance of the C-type (corrosion-resistant in liquid service) cast steels. Additional information on the corrosion resistance of cast steels is contained below and in Corrosion, Volume 13 of ASM Handbook, formerly 9th Edition Metals Handbook. Table 7 Summary of applications for various corrosion-resistant cast steels Alloy CA-15

Characteristics Widely used in mildly corrosive environments; hardenable; good erosion resistance

CA-40

Similar to CA-15 at higher strength level

CA-6NM

Improved properties over CA-15, especially improved resistance to cavitation


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01 Sep 2005

CA-6N

Outstanding combinations of strength, toughness, and weldability with moderately good corrosion resistance

CB-30

Improved performance in oxidizing environments compared to CA-15; excellent resistance to corrosion by nitric acid, alkaline solutions, and many organic chemicals

CB-7Cu-1

Hardenable with good corrosion resistance

CB-7Cu-2

Superior combination of strength, toughness, and weldability with moderately good corrosion resistance

CC-50

Used in highly oxidizing media (hot HNO 3, acid mine waters)

CD-4MCu

Similar to CF-8 in corrosion resistance, but higher strength, hardness, and stress-corrosion cracking resistance; excellent resistance to environments involving abrasion or erosion-corrosion; usefully employed in handling both oxidizing and reducing corrodents

CE-30

Similar to CC-50, but Ni imparts higher strength and toughness levels. A grade available with controlled ferrite

CF-3, CF-8, CF-20, CF-3M, CF-8M, CF-8C, CF-16F

CF types: most widely used corrosion-resistant alloys at ambient and cryogenic temperatures ÃM variations: enhanced resistance to halogen ion and reducing acids ÃC and F variations: used where application does not permit postweld heat treat ÃA grades available with controlled ferrite

CG-8M

Greater resistance to pitting and corrosion in reducing media than CF-8M; not suitable for nitric acids or other strongly oxidizing environments

CH-20

Superior to CF-8 in specialized chemical and paper application in resistance to hot H2SO3, organic acids, and dilute H2SO4; the high nickel and chromium contents also make this alloy less susceptible to intergranular corrosion after exposure to carbide-precipitating temperatures

CK-20

Improved corrosion resistance compared to CH-20

CN-7M

Highly resistant to H2SO4, H3PO4, H2SO3 salts, and seawater. Good resistance to hot chloride salt solutions, nitric acid, and many reducing chemicals

General Corrosion of Martensitic Alloys. The martensitic grades include CA-15, CA-15M, CA-6NM, CA-6NM-B, CA-40, CB-7Cu1, and CB-7Cu-2. These alloys are generally used in applications requiring high strength and some corrosion resistance. Alloy CA-15 typically exhibits a microstructure of martensite and ferrite. This alloy contains the minimum amount of chromium to be considered a stainless steel (11 to 14% Cr) and as such may not be used in aggressive environments. It does, however exhibit good atmospheric-corrosion resistance and it resists staining by many organic environments. Alloy CA-15M may contain slightly more molybdenum than CA-15 (up to 1% Mo) and therefore may have improved general corrosion resistance in relatively mild environments. Alloy CA-6NM is similar to CA-15M except that it contains more nickel and molybdenum, thereby improving its general corrosion resistance. Alloy CA-6NM-B is a lower-carbon version of this alloy. The lower strength level promotes resistance to sulfide stress cracking. Alloy CA-40 is a higher-strength version of CA-15 and it, too, exhibits excellent atmospheric-corrosion resistance after a normalize and temper heat treatment. Microstructurally, the CB-7Cu alloys usually consist of mixed martensite and ferrite and, because of the increased chromium and nickel levels compared to the other martensitic alloys, they offer improved corrosion resistance to seawater and some mild acids. These alloys also have good atmospheric-corrosion resistance. The CB-7Cu alloys are hardenable and offer the possibility of increased strength and improved corrosion resistance among the martensitic alloys. General Corrosion of Ferritic Alloys. Alloys CB-30 and CC-50 are higher-carbon and higher-chromium alloys than are the CA alloys mentioned above. Each alloy is predominantly ferritic, although a small amount of martensite may be found in CB-30. Alloy CB-30 contains 18 to 21% Cr and is used in chemical processing and oil refining applications. The chromium content is sufficient to have good corrosion resistance to many acids, including nitric acid (HNO3). General Corrosion of Austenitic and Duplex Alloys. Alloy CF-8 may be fully austenitic, but it more commonly contains some residual ferrite (3 to 30%) in an austenite matrix. In the solution-annealed condition, this alloy has excellent resistance to a wide variety of acids. It is particularly resistant to highly oxidizing acids, such as boiling HNO3. The duplex nature of the microstructure of this alloy imparts additional resistance to SCC compared to its wholly austenitic counterparts. Alloy CF-3 is a reduced-carbon version of CF-8 with essentially identical corrosion resistance except that CF-3 is much less susceptible to sensitization. For applications in which the corrosion resistance of the weld HAZ may be critical, CF-3 is a common material selection. Alloys CF-8A and CF-3A contain more ferrite than their CF-8 and CF-3 counterparts. Because the higher ferrite content is achieved by increasing the chromium/nickel equivalent ratio, the CF-8A and CF-3A alloys may have slightly higher chromium or slightly lower nickel contents than the low-ferrite equivalents. In general, the corrosion resistance is very similar, but the strength increases with ferrite content. Because of the high ferrite content, service should be restricted to temperatures below 400 °C (750 °F) because of the possibility of severe embrittlement. Alloy CF-8C is the niobium-stabilized grade of the CF-8 alloy class. This alloy contains small amounts of niobium, which tend to form carbides preferentially over chromium carbides and improve


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01 Sep 2005

intergranular corrosion resistance in applications involving relatively high service temperatures. Alloy CF-16F is a selenium-bearing free-machining grade of cast stainless steel. Because CF-16F nominally contains 19% Cr and 10% Ni, it has adequate corrosion resistance to a wide range of corrosive materials but the large number of selenide inclusions makes surface deterioration and pitting definite possibilities. Alloy CE-30 is a nominally 27Cr-9Ni alloy that normally contains 10 to 20% ferrite in an austenite matrix. The high carbon, high ferrite content provides relatively high strength. The high chromium content and duplex structure act to minimize corrosion because of the formation of chromium carbides in the microstructure. This particular alloy is known for good resistance to sulfurous acid and sulfuric acid and is used extensively in the pulp and paper industry (see the article "Corrosion in the Pulp and Paper Industry" in Corrosion, Volume 13 of ASM Handbook, formerly 9th Edition Metals Handbook). Alloy CD-4MCu is the most highly alloyed material in this group of alloys, and a microstructure containing approximately equal amounts of ferrite and austenite is common. The low carbon content and high chromium content render the alloy relatively immune to intergranular corrosion. High chromium and molybdenum provide a high degree of localized corrosion resistance (crevices and pitting), and the duplex microstructure provides SCC resistance in many environments. This alloy can be precipitation hardened to provide strength and is also relatively resistant to abrasion and erosion-corrosion. Fully Austenitic Alloys. Alloys CH-10 and CH-20 are fully austenitic and contain 22 to 26% Cr and 12 to 15% Ni. The high chromium content minimizes the tendency toward the formation of chromium-depleted zones due to sensitization. These alloys are used for handling paper pulp solutions and are known for good resistance to dilute H 2SO4 and HNO3. Alloy CK-20 contains 23 to 27% Cr and 19 to 22% Ni and is less susceptible than CH-20 to intergranular corrosion attack in many acids after brief exposures to the chromium carbide formation temperature range. Maximum corrosion resistance is achieved by solution treatment. Alloy CK-20 possesses good corrosion resistance to many acids and, because of its fully austenitic structure, can be used at relatively high temperatures. Alloy CN-7M exhibits excellent corrosion resistance in a wide variety of environments and is often used for H2SO4 service. Relatively high resistance to intergranular corrosion and SCC make this alloy attractive for many applications. Although CN-7M is relatively highly alloyed, its fully austenitic structure may lead to SCC susceptibility for some environments and stress states. Alloy CF-20 is a fully austenitic, relatively high-strength corrosion-resistant alloy. The 19% Cr content provides resistance to many types of oxidizing acids, but the high carbon content makes it imperative that this alloy be used in the solution-treated condition for environments known to cause intergranular corrosion. Intergranular Corrosion. Ferritic alloys may also be sensitized by the formation of extensive chromium carbide networks, but because of the high bulk chromium content and rapid diffusion rates of chromium in ferrite, the formation of carbides can be tolerated if the alloy has been slowly cooled from a solutionizing temperature of 780 to 900 °C (1435 to 1650 °F). The slow cooling allows replenishment of the chromium adjacent to the carbides. Martensitic alloys normally do not contain sufficient bulk chromium to be used in applications in which intergranular corrosion is likely to be a concern. Austenitic and duplex stainless steels use solution annealing for the prevention or reduction of intergranular corrosion (see "Sensitization and Solution Annealing of Austenitic and Duplex Alloys" in this article). Failure to solution treat a particular alloy or an improper solution treatment may seriously compromise the observed corrosion resistance in service. If solution treatment of the alloy after casting and/or welding is impractical or impossible, the metallurgist has several tools from which to choose to minimize potential intergranular corrosion problems. The low-carbon grades CF-3 and CF-3M are commonly used when heat treatment is impractical or as a solution to the sensitization incurred during welding. The low carbon content, that is, 0.03% C (max), of these alloys precludes the formation of an extensive number of chromium carbides. In addition, these alloys normally contain 3 to 30% ferrite in an austenitic matrix. By virtue of rapid carbide precipitation kinetics at ferrite/austenite interfaces compared to austenite/austenite interfaces, carbide precipitation is confined to ferrite-austenite boundaries in alloys containing a minimum of about 3 to 5% ferrite (Ref 4, 5). If the ferrite network is discontinuous in the austenite matrix (depending on the amount, size, and distribution of ferrite pools), extensive intergranular corrosion will not be a problem in most of the environments to which these alloys will be subjected. The niobium-modified grade of 18-8, known as CF-89C, is produced for similar applications in which heat treatment is impractical. Niobium-containing alloys that have been heated to sensitizing temperatures around 650 °C (1200 °F) are not susceptible to intergranular corrosion. However, they are more susceptible to overall corrosion when tested in nitric acid, compared to the niobium-free, quench-annealed alloys of the same nickel, chromium, and carbon contents. Addition of niobium to molybdenum-containing type CF alloys has also been found unsatisfactory for castings. When both niobium and molybdenum are present, the ferrite phase tends to form as an interconnected network and is especially likely to transform into the brittle σ phase. As a result, castings in the as-cast condition become embrittled and have a tendency to crack. When the niobium-bearing grade CF-8C is in the as-cast condition, most of its carbon is in the form of niobium carbide, precluding chromium carbide precipitation in the critical temperature range from 425 to 870 °C (800 to 1600 °F) and particularly from 565 to 650 °C (1050 to 1200 °F). The alloy CF-8C is solution treated at 1120 °C (2050 °F), quenched to room temperature, and then reheated to 870 to 925 °C (1600 to 1700) °F), at which temperature precipitation of niobium carbide occurs. An alternative method is solution treating at 1120 °C (2050 °F), cooling to the 870 to 925 °C (1600 to 1700 °F) range, and then holding at this temperature before cooling to room temperature. For maximum corrosion resistance, it is recommended that this alloy be solution treated before being stabilized. Weld crack sensitivity of CF alloys containing niobium (CF-8C) is more pronounced in the fully austenitic grade. Cracking may be alleviated through the introduction into the weld deposit of a small amount of ferrite, usually between 4 and 10%.


Page 1422



01 Sep 2005

However, appreciable amounts of ferrite in niobium-bearing corrosion-resistant steels will transform, at least partly, to the σ or χ phase upon heating to between 540 and 925 °C (1000 and 1700 °F). Stress-Corrosion Cracking. The SCC of cast stainless steels has been investigated for only a limited number of environments, heat treatments, and test conditions. From the limited information available, the following generalizations apply. First, SCC resistance seems to improve as the composition is adjusted to provide increasingly greater amounts of ferrite in an austenitic matrix. This trend continues to a certain level, apparently near 50% ferrite (Fig. 6 ). Second, a lower nickel content tends to improve SCC resistance in cast duplex alloys, possibly because of its effect on ferrite content (Ref 6). Third, ferrite appears to be involved in a keying action in discouraging SCC. At low and medium stress levels, the ferrite tends to block the propagation of stress-corrosion cracks. This may be due to a change in composition and/or crystal structure across the austenite/ferrite boundary. As the stress level increases, crack propagation may change from austenite/ferrite boundaries to transgranular propagation (Ref 6, 7). Finally, reducing the carbon content of cast stainless alloys, thereby reducing the susceptibility to sensitization, improves SCC resistance. This is also true for wrought alloys. Fig. 6 Stress required to produce stress-corrosion cracking in several corrosion-resistant cast steels with varying amounts of ferrite

Mechanical Properties of Corrosion-Resistant Cast Steels The importance of mechanical properties in the selection of corrosion-resistant cast steels is established by the casting application. The paramount basis for alloy selection is normally the resistance of the alloy to the specific corrosive media or environment of interest. The mechanical properties of the alloy are usually, but not always, secondary considerations in these applications. Room-Temperature Mechanical Properties. Representative room-temperature tensile properties, hardness, and Charpy impact values for corrosion-resistant cast steels are given in Fig. 7 . These properties are representative of the alloys rather than the specification requirements. Minimum specified mechanical properties for these alloys are given in ASTM standards A 351, A 743, A 744, and A 747. A wide range of mechanical properties are attainable depending on the selection of alloy composition and heat treatment. Tensile strengths ranging from 475 to 1310 MPa (69 to 190 ksi) and hardness from 130 to 400 HB are available among the cast corrosion-resistant alloys. Similarly, wide ranges exist in yield strength, elongation, and impact toughness. Fig. 7 Mechanical properties of various cast corrosion-resistant steels at room temperature. (a) Tensile strength. (b) 0.2% offset yield strength. (b) 0.2% offset yield strength. (c) Charpy keyhole impact energy. (d) Brinell hardness. (e) Elongation. Also given are the heat treatments used for test materials: AC, air cool; FC, furnace cool; WQ, water quench; A, anneal; T, temper.


Page 1423



Alloy

01 Sep 2005

Heat treatment

CA-15 (a)

AC from 980 °C (1800 °F), T at 790 °C (1450 °F)

ÃÃ(b)

AC from 980 °C (1800 °F), T at 650 °C (1200 °F)

ÃÃ(c)

AC from 980 °C (1800 °F), T at 595 °C (1100 °F)

ÃÃ(d)

AC from 980 °C (1800 °F), T at 315 °C (600 °F)

CA-40 (a)

AC from 980 °C (1800 °F), T at 760 °C (1400 °F)

ÃÃ(b)

AC from 980 ° (1800 °F), T at 650 °C (1200 °F)

ÃÃ(c)

AC from 980 °C (1800 °F), T at 595 °C (1100 °F)

ÃÃ(d)

AC from 980 ° (1800 °F), T at 315 °C (600 °F)

CB-30

A at 790 °C (1450 °F), FC to 540 °C (1000 °F), AC

CC-50 (a)

As-cast (2% Ni; >0.15% N)

ÃÃ(c)

AC from 1040 °C (1900 °F) (>2% Ni; >0.15% N)

CE-30 (a)

As-cast

ÃÃ(b)

WQ from 1065−1120 °C (1950−2050 °F)

CF-8

WQ from 1065−1120 °C (1950−2050 °F)


Page 1424



01 Sep 2005

CF-20

WQ from above 1095 °C (2000 °F)

CF-8M, CF-12M

WQ from 1065−1150 °C (1950−2100 °F)

CF-8C

WQ from 1065−1120 °C (1950−2050 °F)

CF-16F


CH-20


CK-20

WQ from above 1150 ° (2100 °F)

CN-7M

WQ from above 1065−1120 °C (1950−2050 °F)

The straight chromium steels (CA-15, CA-40, CB-30, and CC-50) possess either martensitic or ferritic microstructures in the end-use condition (Table 1 ). The CA-15 and CA-40 alloys, which contain nominally 12% Cr, are hardenable through heat treatment by means of the martensite transformation and are of the selected as much or more for their high strength as for their comparatively modest corrosion resistance. The higher-chromium CB-30 and CC-50 alloys, on the other hand, are fully ferritic alloys that are not hardenable by heat treatment. These alloys are generally used in the annealed condition and exhibit moderate tensile properties and hardness. Like most ferritic alloys, CB-30 and CC-50 possess limited impact toughness, especially at low temperatures. Three chromium-nickel alloys, CA-6NM, CB-7Cu, and CD-4M Cu, are exceptional in their response to heat treatment and in the resultant mechanical properties. Alloy CA-6NM is balanced compositionally for martensitic hardening response. This alloy was developed as an alternative to CA-15 and has improved impact toughness and weldability. The CB-7Cu and CD-4MCu alloys both contain copper and can be strengthened by age hardening. These alloys are initially solution heat treated and then cooled rapidly (usually by quenching in oil or water); thus, the phases that would normally precipitate at slow cooling rates cannot form. The casting is then heated to an intermediate aging temperature at which the precipitation reaction can occur under controlled conditions until the desired combination of strength and other properties is achieved. The CB-7Cu alloy possesses a martensitic matrix, while the CD-4MCu alloy possesses a duplex microstructure, consisting of approximately 40% austenite in a ferritic matrix. Alloy CB-7 Cu is applied in the aged condition to obtain the benefit of its excellent combination of strength and corrosion resistance, but alloy CD-4MCu is seldom applied in the aged condition because of its relatively low resistance to SCC in this condition compared to its superior corrosion resistance in the solution-annealed condition. The CE, CF, CG, CH, CN, and CK alloys are essentially not hardenable by heat treatment. To ensure maximum corrosion resistance, however, it is necessary that castings of these grades receive a high-temperature solution anneal (see "Sensitization and Solution Annealing of Austenitic and Duplex Alloys" in this article). By virtue of their microstructures, which are fully austenitic or duplex without significant carbide precipitation, the alloys exhibit generally excellent impact toughness at low temperatures. The tensile strength range represented by these alloys typically extends from 475 to 670 MPa (69 to 97 ksi). As indicated earlier in the section "Austenitic-Ferritic (Duplex) Alloys" in this article, the alloys with duplex structures can be strengthened by balancing the composition for higher ferrite levels (Fig. 5 ). The tensile and yield strengths of CF alloys with a ferrite number of 35 are typically 150 MPa (22 ksi) higher than those of fully austenitic alloys. Tensile ductility (Table 4 ) and impact toughness (Table 5 ) are lowered with increasing ferrite content. Effects from High Temperatures. Cast corrosion-resistant high-alloy steels are used extensively at moderately elevated temperatures (up to 650 °C, or 1200 °F). Elevated-temperature properties are important selection criteria for these applications. Table 8 gives the tensile properties of a corrosion-resistant cast steel at various test temperatures. In addition, mechanical properties after long-term exposure at elevated temperatures are increasingly considered because of the aging effect that these exposures may have. For example, cast alloys CF-8C, CF-8M, CE-30A, and CA-15 are currently used in high-pressure service at temperatures up to 540 °C (1000 °F) in sulfurous acid environments in the petro-chemical industry. Other uses are in the power-generating industry at temperatures up to 565 °C (1050 °F). Table 8 Short-time tensile properties of peripheral-welded cylinders of CF-8 alloy Cylinders were 38 mm (11=2in.) thick; specimens were machined with longitudinal axes perpendicular to welded seam and with seam at middle of gage length. Testing Tensile Yield strength at Proportional Modulus of temperature strength 0.2% offset limit(a) elasticity(a) Reduction Elongation Location of °C

°F

MP a

ksi

MPa

ksi

500

72.5

238

34.5

500

72.5

261

37.8

ksi

in area, %

in 50 mm (2 in.), %

...

...

59.0

49

179

26

62.1

58

MPa

106psi

final rupture

...

...

...

186

27

...

GPa

Base metal Keel block(b) Room 315

600

330

47.8

169

24.5

90

13

54.9

33.5

152

22

...

425

800

339

49.2

167

24.2

59

8.5

58.6

37.5

134

19.5

...

540

1000

291

42.2

140

20.3

55

8

60.8

32.5

117

17

...

595

1100

279

40.4

130

18.8

45

6.5

59.1

38

110

16

...

Welded joint


Page 1425


Room


490

71.0

247

35.8

148

21.5

70.8

01 Sep 2005

42

186

27

Base metal

315

600

341

49.5

199

28.8

72

10.5

58.3

15.5

152

22

Base metal

425

800

355

51.5

171

24.8

69

10

46.3

24.5

131

19

Base metal

540

1000

326

47.3

188

27.3

62

9

62.8

23.5

114

16.5

Base metal

595 1100 272 39.4 134 19.5 55 8 70.4 31 107 15.5 Base metal (a) Values of proportional limit and modulus of elasticity at elevated temperatures are apparent values because creep occurs. (b) Separately cast from same heat as cylinders

Room-temperature properties after exposure to elevated service temperatures may differ from those in the as-heat-treated condition because of the microstructural changes that may take place at the service temperature. Microstructural changes in iron-nickel-chromium-(molybdenum) alloys may involve the formation of carbides and such phases as σ, χ, and η (Laves). The extent to which these phases form depends on the composition, as well as the time at elevated temperature. The martensitic alloys CA-15 and CA-6NM are subject to minor changes in mechanical properties and SCC resistance in NaCl and polythionic acid environments upon exposure for 3000 h at up to 565 °C (1050 °F). In CF-type chromium-nickel-(molybdenum) steels, only negligible changes in ferrite content occur during 10,000 h exposure at 400 °C (750 °F) and during 3000 h exposure at 425 °C (800 °F). Carbide precipitation, however, does occur at these temperatures, and noticeable Charpy V-notch energy losses have been reported. Above 425 °C (800 °F), microstructural changes in chromium-nickel-(molybdenum) alloys take place at an increased rate. Carbides and σ phase form rapidly at 650 °C (1200 °F) at the expense of ferrite. Tensile ductility and Charpy V-notch impact energy (Fig. 8 ) are prone to significant losses under these conditions. Density changes, resulting in contraction, have been reported as a result of these high-temperature exposures. Fig. 8 Charpy V-notch impact energy of three corrosion-resistant cast steels at room temperature after aging at 594 °C (1100 °F). Source: Ref 8


Page 1426



01 Sep 2005

Fatigue Properties and Corrosion Fatigue. The resistance of cast stainless steels to fatigue depends on a sizable number of material, design, and environmental factors. For example, design factors of importance include the stress distribution within the casting (residual and applied stresses), the location and severity of stress concentrators (surface integrity), and the environment and service temperatures. Material factors of importance include strength and microstructure. It is generally found that fatigue strength increases with the tensile strength of a material. Both fatigue strength and tensile strength usually increase with decreasing temperature. Under equivalent conditions of stress, stress concentration, and strength, evidence suggests that austenitic materials are less notch sensitive than martensitic or ferritic materials. Corrosion fatigue is highly specific to the environment and alloy. The martensitic materials are degraded the most in both absolute and relative terms. If left to corrode freely in seawater, they have very little resistance to corrosion fatigue. This is remarkable in view of their very high strength and fatigue resistance in air. Properties can be protected if suitable cathodic protection is applied. However, because these materials are susceptible to hydrogen embrittlement, cathodic protection must be carefully applied. Too large a protective potential will lead to catastrophic hydrogen stress cracking. Austenitic materials are also severely degraded in corrosion fatigue strength under conditions conducive to pitting, such as in seawater. However, they are easily cathodically protected without fear of hydrogen embrittlement and perform well in fresh waters. The corrosion fatigue behavior of duplex alloys has not been widely studied.

Heat-Resistant Cast Steels As previously mentioned, castings are classified as heat resistant if they are capable of sustained operation while exposed, either continuously or intermittently, to operating temperatures that result in metal temperatures in excess of 650 °C (1200 °F). Cast steels for this type of service include iron-chromium (straight chromium), iron-chromium-nickel, and iron-nickel-chromium alloys. In applications of heat-resistant alloys, considerations include:


Page 1427



01 Sep 2005

• Resistance to corrosion at elevated temperatures • Stability (resistance to warping, cracking, or thermal fatigue) • Creep strength (resistance to plastic flow) Table 9 briefly compares the various H-type grades of heat-resistant steel castings in terms of general corrosion resistance and creep values. Table 9 General corrosion characteristics of heat-resistant cast steels and typical limiting creep stress values at indicated temperatures Creep test temperature Alloy

Corrosion characteristics

Limiting creep stress (0.0001%/h)

°C

°F

MPa

ksi

HA

Good oxidation resistance to 650 °C (1200 °F); widely used in oil refining industry

650

1200

21.5

3.1

HC

Good sulfur and oxidation resistance up to 1095 °C (2000 °F); minimal mechanical properties; used in applications where strength is not a consideration or for moderate load bearing up to 650 °C (1200 °F)

870

1600

5.15

0.75

HD

Excellent oxidation and sulfur resistance plus weldability

980

1800

6.2

0.9

HE

Higher temperature and sulfur resistance capabilities than HD

980

1800

9.5

1.4

HF

Excellent general corrosion resistance to 815 °C (1500 °F) with moderate mechanical properties

870

1600

27

3.9

HH(a)

High strength; oxidation resistant to 1090 °C (2000 °F); most widely used

980

1800

7.5 (type I) 14.5 (type II)

1.1 (type I) 2.1 (type II)

H1

Improved oxidation resistance compared to HH

980

1800

13

1.9

HK

Because of its high temperature strength, widely used for stressed parts in structural applications up to 1150 °C (2100 °F); offers good resistance to corrosion by hot gases, including sulfur-bearing gases, in both oxidizing and reducing conditions (although HC, HE, and HI are more resistant in oxidizing gases); used in air, ammonia, hydrogen, and molten neutral salts; widely used for tubes and furnace parts

1040

1900

9.5

1.4

HL

Improved sulfur resistance compared to HK; especially useful where excessive scaling must be avoided

980

1800

15

2.2

HN

Very high strength at high temperatures; resistant to oxidizing and reducing flue gases

1040

1900

11

1.6

HP

Resistant to both oxidizing and carburizing atmospheres at high temperatures

980

1800

19

2.8

HP-50WZ

Improved creep rupture strength at 1090 °C (2000 °F) and above compared to HP

1090

2000

4.8

0.7

HT

Widely used in thermal shock applications; corrosion resistant in air, oxidizing and reducing flue gases, carburizing gases, salts, and molten metals; performs satisfactorily up to 1150 °C (2100 °F) in oxidizing atmospheres and up to 1095 °C (2000 °F) in reducing atmospheres, provided that limiting creep stress values are not exceeded

980

1800

14

2.0

HU

Higher hot strength than HT and often selected for its superior corrosion resistance

980

1800

15

2.2

HW

High hot strength and electrical resistivity; performs satisfactorily to 1120 °C (2050 °F) in strongly oxidizing atmospheres and up to 1040 °C (1900 °F) in oxidizing or reducing products of combustion that do not contain sulfur; resistant to some salts and molten metals

980

1800

9.5

1.4

980

1800

11

1.6

HX

Resistant to hot-gas corrosion under cycling conditions without cracking or warping; corrosion resistant in air, carburizing gases, combustion gases, flue gases, hydrogen, molten cyanide, molten lead, and molten neutral salts at temperatures up to 1150 °C (2100 °F) (a) Two grades: type I (ferrite in austenite) and type II (wholly austenitic), per ASTM A 447

Commercial applications of heat-resistant castings include metal treatment furnaces, gas turbines, aircraft engines, military equipment, oil refinery furnaces, cement mill equipment, petrochemical furnaces, chemical process equipment, power plant equipment, steel mill equipment, turbochargers, and equipment used in manufacturing glass and synthetic rubber. Alloys of the iron-chromium and iron-chromium-nickel groups are of the greatest commercial importance.


Page 1428



01 Sep 2005

General Properties General corrosion and creep properties of heat-resistant steel castings are compared in Table 8 . The compositions of these heat-resistant cast steels are given in Table 2 . These heat-resistant cast steels resemble corrosion-resistant cast steels (Table 1 ) except for their higher carbon content, which imparts greater strength at elevated temperatures. Typical tensile properties of heat-resistant cast steels at room temperature are given in Table 10 and (at elevated temperatures) in Table 11 . Table 10 Typical room-temperature properties of ACI heat-resistant casting alloys Tensile strength

Yield strength

MPa

ksi

MPa

ksi

Elongation, %

Hardness, HB

As-cast

760

110

515

75

19

223

Aged(a)

790

115

550

80

18

...

HD

As-cast

585

85

330

48

16

90

HE

As-cast

655

95

310

45

20

200

Aged(a)

620

90

380

55

10

270

HF

As-cast

635

92

310

45

38

165

Aged(a)

690

100

345

50

25

190

HH, type 1

As-cast

585

85

345

50

25

185

Aged(a)

595

86

380

55

11

200

HH, type 2

As-cast

550

80

275

40

15

180

Aged(a)

635

92

310

45

8

200

As-cast

550

80

310

45

12

180

Aged(a)

620

90

450

65

6

200

As-cast

515

75

345

50

17

170

Aged(b)

585

85

345

50

10

190

Alloy

Condition

HC

HI HK HL

As-cast

565

82

360

52

19

192

HN

As-cast

470

68

260

38

13

160

HP

As-cast

490

71

275

40

11

170

HT

As-cast

485

70

275

40

10

180

Aged(b)

515

75

310

45

5

200

HU

As-cast

485

70

275

40

9

170

Aged(c)

505

73

295

43

5

190

As-cast

470

68

250

36

4

185

Aged(d)

580

84

360

52

4

205

As-cast

450

65

250

36

9

176

HW HX

Aged(c) 505 73 305 44 9 185 (a) Aging treatment: 24 h at 760 °C (1400 °F), furnace cool. (b) Aging treatment: 24 h at 760 °C (1400 °F), air cool. (c) Aging treatment: 48 h at 980 °C (1800 °F), air cool. (d) Aging treatment: 48 h at 980 °C (1800 °F), furnace cool

Table 11 Representative short-term tensile properties of cast heat-resistant alloys at elevated temperatures Property at indicated temperature 760 °C (1400 °F) Ultimate tensile strength ksi

870 °C (1600 °F)

Yield strength Ultimate at tensile Elonga 0.2% offset strength tion, % MPa ksi MPa ksi

Alloy

MPa

HA

462(a 67(a) 220(b) 32(b) )

980 °C (1800 °F)

Yield Ultimate strength at Elongati tensile 0.2% offset strength on, % MPa ksi MPa ksi

Yield strength at Elongati 0.2% offset on, % MPa ksi

...

...

...

...

...

...

...

...

...

...

...

HD

248

36

...

...

14

159

23

...

...

18

103

15

...

...

40

HF

262

38

172

25

16

145

21

107

15.5

16

...

...

...

...

...

HH (type I)(c)

228

33

117

17

18

127

18.5

93

13.5

30

62

9

43

6.3

45


Page 1429



01 Sep 2005

HH (type II)(c)

258

37.4

136

19.8

16

148

21.5

110

16

18

75

10.9

50

7.3

31

HI

262

38

...

...

6

179

26

...

...

12

...

...

...

...

...

HK

258

37.5

168

24.4

12

161

23

101

15

16

85.5

12.4

60

8.7

42

HL

345

50

...

...

...

210

30.5

...

...

...

129

18.7

...

...

...

HN

...

...

...

...

...

140

20

100

14.5

37

83

12

66

9.6

51

HP

296

43

200

29

15

179

26

121

17.5

27

100

14.5

76

11

46

HT

240

35

180

26

10

130

19

103

15

24

76

11

55

8

28

HU

275

40

...

...

...

135

19.5

...

...

20

69

10

43

6.2

28

HW

220

32

158

23

...

131

19

103

15

...

69

10

55

8

40

HX

310(d 45(d) 138(d) 20(d) 8(d) 141 20.5 121 17.5 48 74 10.7 4.7 6.9 40 ) (a) In this instance, test temperature was 540 °C (1000 °F). (b) Test temperature was 590 °C (1100 °F). (c) Type I and II per ASTM A 447 (d) Test temperature was 650 °C (1200 °F).

Iron-chromium alloys contain 10 to 30% Cr and little or no nickel. These alloys are useful chiefly for resistance to oxidation; they have low strength at elevated temperatures. Use of these alloys is restricted to conditions, either oxidizing or reducing, that involve low static loads and uniform heating. Chromium content depends on anticipated service temperature. Iron-chromium-nickel alloys contain more than 13% Cr and more than 7% Ni (always more chromium than nickel). These austenitic alloys are ordinarily used under oxidizing or reducing conditions similar to those withstood by the ferritic iron-chromium alloys, but in service they have greater strength and ductility than the straight chromium alloys. They are used, therefore, to withstand greater loads and moderate changes of temperature. These alloys also are used in the presence of oxidizing and reducing gases that are high in sulfur content. Iron-nickel-chromium alloys contain more than 25% Ni and more than 10% Cr (always more nickel than chromium). These austenitic alloys are used for withstanding reduction as well as oxidizing atmospheres, except where sulfur content is appreciable. (In atmospheres containing 0.05% or more hydrogen sulfide, for example, iron-chromium-nickel alloys are recommended.) In contrast with iron-chromium-nickel alloys, iron-nickel-chromium alloys do not carburize rapidly or become brittle and do not take up nitrogen in nitriding atmospheres. These characteristics become enhanced as nickel content is increased, and in carburizing and nitriding atmospheres casting life increases with nickel content. Austenitic iron-nickel-chromium alloys are used extensively under conditions of severe temperature fluctuations such as those encountered by fixtures used in quenching and by parts that are not heated uniformly or that are heated and cooled intermittently. In addition, these alloys have characteristics that make them suitable for electrical resistance heating elements. Metallurgical Structures The structures of chromium-nickel and nickel-chromium cast steels must be wholly austenitic, or mostly austenitic with some ferrite, if these alloys are to be used for heat-resistant service. Depending on the chromium and nickel content (see the section "Composition and Microstructure" in this article), the structures of these iron-base alloys can be austenitic (stable), ferritic (stable, but also soft, weak, and ductile) or martensitic (unstable); therefore, chromium and nickel levels should be selected to achieve good strength at elevated temperatures combined with resistance to carburization and hot-gas corrosion. A fine dispersion of carbides or intermetallic compounds in an austenitic matrix increases high-temperature strength considerably. For this reason, heat-resistant cast steels are higher in carbon content than are corrosion-resistant alloys of comparable chromium and nickel content. By holding at temperatures where carbon diffusion is rapid (such as above 1200 °C) and then rapidly cooling, a high and uniform carbon content is established, and up to about 0.20% C is retained in the austenite. Some chromium carbides are present in the structures of alloys with carbon contents greater than 0.20%, regardless of solution treatment, as described in the section "Sensitization and Solution Annealing of Austenitic and Duplex Alloys" in this article. Castings develop considerable segregation as they freeze. In standard grades, either in the as-cast condition or after rapid cooling from a temperature near the melting point, much of the carbon is in supersaturated solid solution. Subsequent reheating precipitates excess carbides. The lower the reheating temperature, the slower the reaction and the finer the precipitated carbides. Fine carbides increase creep strength and decrease ductility. Intermetallic compounds such as Ni3Al, if present, have a similar effect. Reheating material containing precipitated carbides in the range between 980 and 1200 °C (1800 and 2200 °F) will agglomerate and spheroidize the carbides, which reduces creep strength and increases ductility. Above 1100 °C (2000 °F), so many of the fine carbides are dissolved or spheroidized that this strengthening mechanism loses its importance. For service above 1100 °C (2000 °F), certain proprietary alloys of the iron-nickel-chromium type have been developed. Alloys for this service contain tungsten to form tungsten carbides, which are more stable than chromium carbides at these temperatures. Aging at a low temperature, such as 760 °C (1400 °F), where a fine, uniformly dispersed carbide precipitate will form, confers a high level of strength that is retained at temperatures up to those at which agglomeration changes the character of the carbide dispersion (overaging temperatures). Solution heat treatment or quench annealing, followed by aging, is the treatment generally employed to attain maximum creep strength. Ductility is usually reduced when strengthening occurs; but in some alloys the strengthening treatment correct an unfavorable


Page 1430



01 Sep 2005

grain-boundary network of brittle carbides, and both properties benefit. However, such treatment is costly and may warp castings excessively. Hence, this treatment is applied to heat-resistant castings only for the small percentage of applications for which the need for prenium performance justifies the high cost. Carbide networks at grain boundaries are generally undesirable in iron-base heat-resistant alloys. Grain-boundary networks usually occur in very-high-carbon alloys or in alloys that have cooled slowly through the high-temperature ranges in which excess carbon in the austenite is rejected as grain-boundary networks rather than as dispersed particles. These networks confer brittleness in proportion to their continuity. Carbide networks also provide paths for selective attack in some atmospheres and in certain molten salts. Therefore, it is advisable in some salt bath applications to sacrifice the high-temperature strength imparted by high carbon content and gain resistance to intergranular corrosion by specifying that carbon content be no greater than 0.08%. Straight Chromium Heat-Resistant Castings Iron-chromium alloys, also known as straight chromium alloys, contain either 9 or 28% Cr. HC and HD alloys are included among the straight chromium alloys, although they contain low levels of nickel. HA alloy (9Cr-1Mo), a heat treatable material, contains enough chromium to provide good resistance to oxidation at temperatures up to about 650 °C (1200 °F). The 1% molybdenum is present to provide increased strength. HA alloy castings are widely used in oil refinery service. A higher-chromium modification of this alloy (12 to 14% Cr) is widely used in the glass industry. HA alloy has a structure that is essentially ferritic; carbides are present in pearlitic areas or as agglomerated particles, depending on prior heat treatment. Hardening of the alloy occurs upon cooling in air from temperatures above 815 °C (1500 °F). In the normalized and tempered condition, the alloy exhibits satisfactory toughness throughout its useful temperature range. HC alloy (28% Cr) resists oxidation and the effects of high-sulfur flue gases at temperatures up to 1100 °C (2000 °F). It is used for applications in which strength is not a consideration, or in which only moderate loads are involved, at temperatures of about 650 °C (1200 °F). It is also used where appreciable nickel cannot be tolerated, as in very-high-sulfur atmospheres, or where nickel may act as an undesirable catalyst and destroy hydrocarbons by causing them to crack. HC alloy is ferritic at all temperatures. Its ductility and impact strength are very low at room temperature and its creep strength is very low at elevated temperatures unless some nickel is present. In a variation of HC alloy that contains more than 2% Ni, substantial improvement in all three of these properties is obtained by increasing the nitrogen content to 0.15% or more. HC alloy becomes embrittled when heated for prolonged periods at temperatures between 400 and 550 °C (750 and 1025 °F), and it shows low resistance to impact. The alloys is magnetic and has a low coefficient of thermal expansion, comparable to that of carbon steel. It has about eight times the electrical resistivity and about half the thermal conductivity of carbon steel. Its thermal conductivity, however, is roughly double the value for austenitic iron-chromium-nickel alloys. HD alloy (28Cr-5Ni) is very similar in general properties to HC, except that its nickel content gives it somewhat greater strength at high temperatures. The high chromium content of this alloy makes it suitable for use in high-sulfur atmospheres. HD alloy has a two-phase, ferrite-plus-austenite structure that is not hardenable by conventional heat treatment. Long exposure at 700 to 900 °C (1300 to 1650 °F), however, may result in considerable hardening and severe loss of room-temperature ductility through the formation of σ phase. Ductility may be restored by heating uniformly to 980 °C (1800 °F) or higher and then cooling rapidly to below 650 °C (1200 °F). Iron-Chromium-Nickel Heat-Resistant Castings Heat-resistant ferrous alloys in which the chromium content exceeds the nickel content are made in compositions ranging from 20Cr-10Ni to 30Cr-20Ni. HE alloy (28Cr-10Ni) has excellent resistance to corrosion at elevated temperatures. Because of its higher chromium content, it can be used at higher temperatures than HF alloy and is suitable for applications up to 1100 °C (2000 °F). This alloy is stronger and more ductile at room temperature than the straight chromium alloys. In the as-cast condition, HE alloy has a two-phase, austenite-plus-ferrite structure containing carbides. HE castings cannot be hardened by heat treatment; however, as with HD castings, long exposure to temperatures near 815 °C (1500 °F) will promote formation of σ phase and consequent embrittlement of the alloy at room temperature. The ductility of this alloy can be improved somewhat by quenching from about 1100 °C (2000 °F). Castings of HE alloy have good machining and welding properties. Thermal expansion is about 50% greater than that of either carbon steel or the Fe-Cr alloy HC. Thermal conductivity is much lower than for HD or HC, but electrical resistivity is about the same. HE alloy is weakly magnetic. HE alloy (20Cr-10Ni) is the cast version of 18-8 stainless steel, which is widely used for its outstanding resistance to corrosion. HF alloy is suitable for use at temperatures up to 870 °C (1600 °F). When this alloy is used for resistance to oxidation at elevated temperatures, it is not necessary to keep the carbon content at the low level specified for corrosion-resistant castings. Molybdenum, tungsten, niobium, and titanium are sometimes added to the basic HF composition to improve elevated-temperature strength. In the as-cast condition, HF alloy has an austenitic matrix that contains interdendritic eutectic carbides and, occasionally, a lamellar constituent presumed to consist of alternating platelets of austenite and carbide or carbonitride. Exposure at service


Page 1431



01 Sep 2005

temperatures usually promotes precipitation of finely dispersed carbides, which increases room-temperature strength and causes some loss of ductility. If improperly balanced, as-cast HF may be partly ferritic. HF is susceptible to embrittlement due to σ-phase formation after long exposure at 760 to 815 °C (1400 to 1500 °F). HH Alloy (26Cr-12Ni). Alloys of this nominal composition comprise about one-third of the total production of iron-base heat-resistant castings. Alloy HH is basically austenitic and holds considerable carbon in solid solution, but carbides, ferrite (soft, ductile, and magnetic) and σ (hard, brittle, and nonmagnetic) may also be present in the microstructure. The amounts of the various structural constituents present depend on composition and thermal history. In fact, two distinct grades of material can be obtained within the stated chemical compositional range of the type alloy HH. These grades are defined as type I (partially ferritic) and type II (wholly austenitic) in ASTM A 447. The partially ferritic (type I) alloy HH is adapted to operating conditions that are subject to changes in temperature level and applied stress. A plastic extension in the weaker, ductile ferrite under changing load tends to occur more readily than in the stronger austenitic phase, thereby reducing unit stresses and stress concentrations and permitting rapid adjustment to suddenly applied overloads without cracking. Near 870 °C (1600 °F), the partially ferritic alloys tend to embrittle from the development of σ phase, while close to 760 °C (1400 °F), carbide precipitation may cause comparable loss of ductility. Such possible embrittlement suggests that 930 to 1090 °C (1700 to 2000 °F) is the best service temperature range, but this is not critical for steady temperature conditions in the absence of unusual thermal or mechanical stresses. To achieve maximum strength at elevated temperatures, the HH alloy must be wholly austenitic. Where load and temperature conditions are comparatively constant, the wholly austenitic (type II) alloy HH provides the highest creep strength and permits the use of maximum design stress. The stable austenitic alloy is also favored for cyclic temperature service that might induce σ-phase formation in the partially ferritic type. When HH alloy is heated to between 650 and 870 °C (1200 and 1600 °F), a loss in ductility may be produced by either of two changes within the alloy: precipitation of carbides or transformation of ferrite to σ. When the composition is balanced so that the structure is wholly austenitic, only carbide precipitation normally occurs. In partly ferritic alloys, both carbides and σ phase may form. The wholly austenitic (type II) HH alloy is used extensively in high-temperature applications because of its combination of relatively high strength and oxidation resistance at temperatures up to 1100 °C (2000 °F). Typical tensile properties and impact toughness of the type II HH alloy at elevated temperatures are shown in Fig. 9 (a). The HH alloy (type I or II) is seldom used for carburizing applications because of embrittlement from carbon absorption. High silicon content (over 1.5%) will fortify the alloy against carburization under mild conditions, but will promote ferrite formation and possible σ embrittlement. Fig. 9 Effect of short-term elevated-temperature exposure on the tensile properties of wholly austenitic (type II) HH cast steel (a) and of five other heat-resistant cast steels: (b) HF cast steel, (c) HK-40 cast steel, (d) HN cast steel, (e) HP cast steel, and (f) HT cast steel. Long-term elevated-temperature exposure reduces the strengthening effects between 500 to 750 °C (900 to 1400 °F) in (c), (d), and (e). Tensile properties of alloy HT in (f) include extrapolated data (dotted lines) below 750 °C but should be similar to alloy HN in terms of yield and tensile strengths. Source: Ref 9


Page 1432



01 Sep 2005

For the wholly austenitic (type II) HH alloy, composition balance is critical in achieving the desired austenitic microstructure (see "Composition and Microstructure" in this article). An imbalance of higher levels of ferrite-promoting elements compared to levels of austenite-promoting elements may result in substantial amounts of ferrite which improves ductility, but decreases strength at high temperatures. If a balance is maintained between ferrite-promoting elements (such as chromium and silicon) and austenite-promoting elements (such as nickel, carbon, and nitrogen), the desired austenitic structure can be obtained. In commercial HH alloy castings, with the usual carbon, nitrogen, manganese, and silicon contents, the ratio of chromium to nickel necessary for a stable austenitic structure is expressed by the inequality:

(Eq 3) Silicon and molybdenum have definite effects on the formation of σ phase. A silicon content in excess of 1% is equivalent to a chromium content three times as great, and any molybdenum content is equivalent to a chromium content four times as great. Before HH alloy is selected as a material for heat-resistant castings, it is advisable to consider the relationship between chemical composition and operating-temperature range. For castings that are to be exposed continuously at temperatures appreciably above 870 °C (1600 °F), there is little danger of severe embrittlement from either the precipitation of carbide or the formation of σ phase, and composition should be 0.50% C (max) (0.35 to 0.40% preferred), 10 to 12% Ni, and 24 to 27% Cr. On the other hand, castings to be used at temperatures from 650 to 870 °C (1200 to 1600 °F) should have compositions of 0.40% C (max), 11 to 14% Ni, and 23 to 27% Cr. For applications involving either of these temperature ranges, that is, 650 to 870 °C (1200 to 1600 °F), or appreciably above 870 °C (1600 °F), composition should be balanced to provide an austenitic structure. For service from 650 to 870 °C (1200 to 1600 °F), for example, a combination of 11% Ni and 27% Cr is likely to produce σ phase and its associated embrittlement, which occurs most rapidly around 870 °C (1600 °F). It is preferable, therefore, to avoid using the maximum chromium content with the minimum nickel content. Short-time tensile testing of fully austenitic HH alloys shows that tensile strength and elongation depend on carbon and nitrogen contents. For maximum creep strength, HH alloy should be fully austenitic in structure (Fig. 10 ). In design of load-carrying castings, data concerning creep stresses should be used with an understanding of the limitations of such data. An extrapolated limiting creep stress for 1% elongation in 10,000 h cannot necessarily be sustained for that length of time without structural damage. Stress-rupture testing is a valuable adjunct to creep testing and a useful aid in selecting section sizes to obtain appropriate levels of design stress. Fig. 10 Creep strength of heat-resistant alloy castings (HT curve is included in both graphs for ease of comparison). Source: Ref 10


Page 1433



01 Sep 2005

Because HH alloys of wholly austenitic structure have greater strength at high temperatures than partly ferritic alloys of similar composition, measurement of ferrite content is recommended. Although a ratio calculated from Eq 3 that is less than 1.7 indicates wholly austenitic material, ratios greater than 1.7 do not constitute quantitative indications of ferrite content. It is possible, however, to measure ferrite content by magnetic analysis after quenching from about 1100 °C (2000 °F). The magnetic permeability of HH alloys increases with ferrite content. This measurement of magnetic permeability, preferably after holding 24 h at 1100 °C (2000 °F) and then quenching in water, can be related to creep strength, which also depends on structure. HH alloys are often evaluated by measuring percentage elongation in room-temperature tension testing of specimens that have been held 24 h at 760 °C (1400 °F). Such a test may be misleading because there is a natural tendency for engineers to favor compositions that exhibit the greatest elongation after this particular heat treatment. High ductility values are often measured for alloys that have low creep resistance, but, conversely, low ductility values do not necessarily connote high creep resistance. HI alloy (28Cr-15Ni) is similar to HH but contains more nickel and chromium. The higher chromium content makes HI more resistant to oxidation than HH, and the additional nickel serves to maintain good strength at high temperatures. Exhibiting adequate strength, ductility, and corrosion resistance, this alloy has been used extensively for retorts operating with an internal


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vacuum at a continuous temperature of 1175 °C (2150 °F). It has an essentially austenitic structure that contains carbides and that, depending on the exact composition balance, may or may not contain small amounts of ferrite. Service at 760 to 870 °C (1400 to 1600 °F) results in precipitation of finely dispersed carbides, which increases strength and decreases ductility at room temperature. At service temperatures above 1100 °C (2000 °F), however, carbides remain in solution, and room-temperature ductility is not impaired. HY alloy (26Cr-20Ni) is somewhat similar to wholly austenitic HH alloy in general characteristics and mechanical properties. Although less resistant to oxidizing gases than HC, HE or HI (Table 12 ), HK alloy contains enough chromium to ensure good resistance to corrosion by hot gases, including sulfur-bearing gases, under both oxidizing and reducing conditions. The high nickel content of this alloy helps make it one of the strongest heat-resistant casting alloys at temperatures above 1040 °C (1900 °F). Accordingly, HK alloy castings are widely used for stressed parts in structural applications at temperatures up to 1150 °C (2100 °F). As normally produced, HK is a stable austenitic alloy over its entire range of service temperatures. The as-cast microstructure consists of an austenitic matrix containing relatively large carbides in the forms of either scattered islands or networks. After the alloy has been exposed to service temperature, fine, granular carbides precipitate within the grains of austenite and, if the temperature is high enough, undergo subsequent agglomeration. These fine, dispersed carbides contribute to creep strength. A lamellar constituent that resembles pearlite, but that is presumed to be carbide or carbonitride platelets in austenite, is also frequently observed in HK alloy. Table 12 Approximate rates of corrosion for ACI heat-resistant casting alloys in air and in flue gas Corrosion rate, mm/yr, at 980 °C (1800 °F) in flue gas with sulfur content of:

Oxidation rate in air, mm/yr

0.12 g/m3

2.3 g/m3

870 °C (1600 °F)

980 °C (1800 °F)

1090 °C (2000 °F)

Oxidizing

Reducing

Oxidizing

Reducing

HB

0.63−

6.25−

12.5−

2.5+

12.5

6.25−

12.5

HC

0.25

1.25

1.25

0.63−

0.63+

0.63

0.63−

HD

0.25−

1.25−

1.25−

0.63−

0.63−

0.63−

0.63−

HE

0.13−

0.63−

0.88−

0.63−

0.63−

0.63−

0.63−

HF

0.13−

1.25+

2.5

1.25+

2.5+

1.25+

6.25

HH

0.13−

0.63−

1.25

0.63−

0.63

0.63

0.63−

HI

0.63−

0.63−

0.63−

0.63−

Alloy

0.13−

0.25+

0.88−

HK

0.25−

0.25−

0.88−

0.63−

0.63−

0.63−

0.63−

HL

0.25+

0.63−

0.88

0.63−

0.63−

0.63

0.63−

HN

0.13

0.25+

1.25−

0.63−

0.63−

0.63

0.63

HP

0.63−

0.63

1.25

0.63−

0.63−

0.63−

0.63−

HT

0.13−

0.25+

1.25

0.63

0.63−

0.63

2.5

HU

0.13−

0.25−

0.88−

0.63−

0.63−

0.63−

0.63

HW

0.13−

0.25−

0.88

0.63

0.63−

1.25−

6.25

0.13−

0.25−

0.88−

0.63−

0.63−

0.63−

0.63−

HX Source: Ref 11

Unbalanced compositions are possible within the standard composition range for HK alloy, and hence some ferrite may be present in the austenitic matrix. Ferrite will transform to brittle σ phase if the alloy is held for more than a short time at about 815 °C (1500 °F), with consequent embrittlement upon cooling to room temperature. Direct transformation of austenite to σ phase can occur in HK alloy in the range of 760 to 870 °C (1400 to 1600 °F), particularly at lower carbon levels (0.20 to 0.30%). The presence of σ phase can cause considerable scatter in property values at intermediate temperatures. The minimum creep rate and average rupture life of HK are strongly influenced by variations in carbon content. Under the same conditions of temperature and load, alloys with higher carbon content have lower creep rates and longer lives than lower-carbon compositions. Room-temperature properties after aging at elevated temperatures are affected also: The higher the carbon, the lower the residual ductility. For these reason, three grades of HK alloys with carbon ranges narrower than the standard HK alloy in Table 2 are recognized: HK-30, HK-40, and HK-50. In these designations, the number indicates the midpoint of a 0.10% C range. HK-40 (Table 2 ) is widely used for high-temperature processing equipment in the petroleum and petro-chemical industries. Figure 9 (c) shows the effect of short-term temperature exposure on an HK-40 alloy. Figure 11 indicates the statistical spread in room-temperature mechanical properties obtained for an HK alloy. These data were obtained in a single foundry and are based on 183 heats of the same alloy. Fig. 11 Statistical spread in mechanical properties of HK alloy. Data are for 183 heats of HK alloy produced in a single


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foundry. Tests were performed at room temperature on as-cast material.

HL alloy (30Cr-20Ni) is similar to HK; its higher chromium content gives it greater resistance to corrosion by hot gases, particularly those containing appreciable amounts of sulfur. Because essentially equivalent high-temperature strength can be obtained with either HK or HL, the superior corrosion resistance of HL makes it especially useful for service in which excessive scaling must be avoided. The as-cast and aged microstructures of HL alloy, as well as its physical properties and fabricating characteristics, are similar to those of HK. Iron-Nickel-Chromium Heat-Resistant Castings Iron-nickel-chromium alloys generally have more stable structures than those of iron-base alloys in which chromium is the predominant alloying element. There is no evidence of an embrittling phase change in iron-nickel-chromium alloys that would impair their ability to withstand prolonged service at elevated temperature. Experimental data indicate that composition limits are not critical; therefore, the production of castings from these alloys does not require the close composition control necessary for making castings from iron-chromium-nickel alloys. The following general observations should be considered in the selection of iron-nickel-chromium alloys: • As nickel content is increased, the ability of the alloy to absorb carbon from a carburizing atmosphere decreases • As nickel content is increased, tensile strength at elevated temperatures decreases somewhat, but resistance to thermal shock and thermal fatigue increases • As chromium content is increased, resistance to oxidation and to corrosion in chemical environments increases • As carbon content is increased, tensile strength at elevated temperatures increases • As silicon content is increased, tensile strength at elevated temperatures decreases, but resistance to carburization increases somewhat HN alloy (25Ni-20Cr) contains enough chromium for good high-temperature corrosion resistance. HN has mechanical properties somewhat similar to those of the much more widely used HT alloy, but has better ductility (see Fig. 9 d and 9 f for a comparison of HN and HT tensile properties above 750 °C, or 1400 °F). It is used for highly stressed components in the temperature range of 980 to 1100 °C (1800 to 2000 °F). In several specialized applications (notably, brazing fixtures), it has given satisfactory service at temperatures from 1100 to 1150 °C (2000 to 2100 °F). HN alloy is austenitic at all temperatures: Its composition limits lie well within the stable austenite field. In the as-cast condition it contains carbide areas, and additional fine carbides precipitate with aging. HN is not susceptible to σ phase formation, and increases in its carbon content are not especially detrimental to ductility. HP, HT, HU, HW, and HX alloys make up about one-third of the total production of heat-resistant alloy castings. When used for fixtures and trays for heat treating furnaces, which are subjected to rapid heating and cooling, these five high-nickel alloys have exhibited excellent service life. Because these compositions are not as readily carburized as iron-chromium-nickel alloys, they are used extensively for parts of carburizing furnaces. Because they form an adherent scale that does not flake off, castings of these alloys are also used in enameling applications in which loose scale would be detrimental. Four of these high-nickel alloys (HT, HU, HW, and HX) also exhibit good corrosion resistance with molten salts and metal. They have excellent corrosion resistance to tempering and to cyaniding salts and fair resistance to neutral salts, with proper control. With molten metal, these alloy exhibit excellent resistance to molten lead, good resistance to molten tin to 345 °C (650 °F), and good resistance to molten cadmium to 410 °C (775 °F). The alloys have poor resistance to antimony, babbitt, soft solder, and similar metal. In many respects, there are no sharp lines of demarcation among the HP, HT, HU, HW, and HX alloys with respect to service applications.


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HP alloy (35Ni-26Cr) is related to HN and HT alloys, but is higher in alloy content. It contains the same amount of chromium but more nickel than HK, and the same amount of nickel but more chromium than HT. This combination of elements makes HP resistant to both oxidizing and carburizing atmospheres at high temperatures. It has creep-rupture properties that are comparable to, or better than, those of HK-40 and HN alloys (Fig. 12 ). Fig. 12 Stress-rupture properties of several heat-resistant alloy castings. (a) 10,000 h rupture stress. (b) 100,000 h rupture stress. Source: Ref 10

HP alloy is austenitic at all temperatures, and is not susceptible to σ-phase formation. Its microstructure consists of massive


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primary carbides in an austenitic matrix; in addition, fine secondary carbides are precipitated within the austenite grains upon exposure to elevated temperatures. This precipitation of carbides is responsible for the strengthening between 500 and 750 °C (900 and 1400 °F) in Fig.(e) 9 . This strengthening, which is reduced after long-term exposure at high temperatures, also occurs for the cast stainless steels shown in Fig. (c) 9 and (d) 9 . HT alloy (35Ni-17Cr) contains nearly equal amounts of iron and alloying elements. Its high nickel content enables it to resist the thermal shock of rapid heating and cooling. In addition, HT is resistant to high-temperature oxidation and carburization and has good strength at the temperatures ordinarily used for heat treating steel. Except in high-sulfur gases, and provided that limiting creep-stress values are not exceeded, it performs satisfactorily in oxidizing atmospheres at temperatures up to 1150 °C (2100 °F) and in reducing atmospheres at temperatures up to 1100 °C (2000 °F). HT alloy is widely used for highly stressed parts in general heat-resistant applications. It has an austenitic structure containing carbides in amounts that vary with carbon content and thermal history. In the as-cast condition, it has large carbide areas at interdendritic boundaries; but fine carbides precipitate within the grains after exposure to service temperature, causing a decrease in room-temperature ductility. Increases in carbon content may decrease the high-temperature ductility of the alloy. A silicon content above about 1.6% provides additional protection against carburization, but at some sacrifice in elevated-temperature strength. HT can be made still more resistant to thermal shock by the addition of up to 2% niobium. HU alloy (39Ni-18Cr) is similar to HT, but its higher chromium and nickel contents give it greater resistance to corrosion by either oxidizing or reducing hot gases, including those that contain sulfur in amounts up to 2.3 g/m 3 (see Table 12 ). Its high-temperature strength and resistance to carburization are essentially the same as those of HT and thus its superior corrosion resistance makes it especially well suited for severe service involving high stress and/or rapid thermal cycling, in combination with an aggressive environment. HW alloy (60Ni-12Cr) is especially well suited for applications in which wide and/or rapid fluctuations in temperature are encountered. In addition, HW exhibits excellent resistance to carburization and high-temperature oxidation. HW alloy has good strength at steel-treating temperatures, although it is not as strong as HT. HW performs satisfactorily at temperatures up to about 1120 °C (2050 °F) in strongly oxidizing atmospheres and up to 1040 °C (1900 °F) in oxidizing or reducing products of combustion, provided that sulfur is not present in the gas. The generally adherent nature of its oxide scale makes HW suitable for enameling furnace service, where even small flakes of dislodged scale could ruin the work in process. HW alloy is widely used for intricate heat-treating fixtures that are quenched with the load and for many other applications (such as furnace retorts and muffles) that involve thermal shock, steep temperature gradients, and high stresses. Its structure is austenitic and contains carbides in amounts that vary with carbon content and thermal history. In the as-cast condition, the microstructure consists of a continuous interdendritic network of elongated eutectic carbides. Upon prolonged exposure at service temperatures, the austenitic matrix becomes uniformly peppered with small carbide particles except in the immediate vicinity of eutectic carbides. This change in structure is accompanied by an increase in room-temperature strength, but there is no change in ductility. HX alloy (66Ni-17Cr) is similar to HW, but contains more nickel and chromium. Its higher chromimum content gives it substantially better resistance to corrosion by hot gases (ever sulfur-bearing gases), which permits it to be used in severe service applications at temperatures up to 1150 °C (2100 °F). However, it has been reported that HX alloy decarburized rapidly at temperatures from 1100 to 1150 °C (2000 to 2100 °F). High-temperatures strength (Table 11 ), resistance to thermal fatigue, and resistance to carburization are essentially the same as for HW; hence HX is suitable for the same general applications in which corrosion must be minimized. The as-cast and aged microstructure of HX, as well as its mechanical properties and fabricating characteristics, are similar to those of HW. Properties of Heat-Resistant Alloys Elevated-Temperature Tensile Properties. The short-term elevated-temperature test, in which a standard tension test bar is heated to a designated uniform temperature and then strained to fracture at a standardized rate, identifies the stress due to a short-term overload that will cause fracture in uniaxial loading. The manner in which the values of tensile strength and ductility change with increasing temperature is shown in Fig. 9 for selected alloys. Representative tensile properties at temperatures between 760 and 980 °C (1400 and 1800 °F) are given in Table 11 for several heat-resistant cast steel grades. Creep and Stress-Rupture Properties. Creep is defined as the time-dependent strain that occurs under load at elevated temperature and is operative in most applications of heat-resistant high-alloy castings at the normal service temperatures. In time, creep may lead to excessive deformation and even fracture at stress considerably below those determined in room-temperature and elevated-temperature short-term tension tests. When the rate or degree of deformation is the limiting factor, the design stress is based on the minimum creep rate and design life after allowing for initial transient creep. The stress that produces a specified minimum creep rate of an alloy or a specified amount of creep deformation in a given time (for example, 1% total creep in 100,000 h) is referred to as the limiting creep strength, or limiting stress. Table 9 lists the creep strength of various H-type castings at specific temperatures. Figure 10 shows creep rates as a function of temperature. Stress-rupture testing is a valuable adjunct to creep testing and is used to select the section sizes necessary to prevent creep rupture of a component. Figure 12 compares the creep-rupture strength of various H-type steels castings at 10,000 and 100,000 h. It should be recognized that long-term creep and stress-rupture values (for example, 100,000 h) are often extrapolated from shorter-term tests. Whether these property values are extrapolated or determined directly often has little bearing on the operating life of high-temperatures parts. The actual material behavior is often difficult to predict accurately because of the complexity of


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the service stresses relative to the idealized, uniaxial loading conditions in the standardized tests and because of the attenuating factors such as cyclic loading, temperature fluctuations, and metal loss from corrosion. The designer should anticipate the synergistic effects of these variables. Thermal fatigue failure involves cracking caused by heating and cooling cycles. Very little experimental thermal fatigue information is available on which to base a comparison of the various alloys, and no standard test as yet has been adopted. Field experience indicates that resistance to thermal fatigue is usually improved with an increase in nickel content. Niobium-modified alloys have been employed successfully when a high degree of thermal fatigue resistance is desired such as in reformer outlet headers. Thermal Shock Resistance. Thermal shock failure may occur as a result of a single, rapid temperature change or as a result of rapid cyclic temperature changes, which induce stress that are high enough to cause failure. Thermal shock resistance is influenced by the coefficient of thermal expansion and the thermal conductivity of materials. Increases in the thermal expansion coefficient or decreases in thermal conductivity reduce the resistance against thermal shock. Table 13 lists the thermal conductivities and expansion coefficients for heat-resistant castings at various temperatures. The HA, HC, and HD alloys, because of their predominately ferritic microstructure, have the lowest thermal expansion coefficients and the highest thermal conductivities. Table 13 Thermal conductivity and mean coefficient of linear thermal expansion of ACI heat-resistant cast steels at various temperatures Mean coefficient of linear thermal expansion for a temperature change

Alloy

Thermal conductivity, W/m · K, at:

From 21 to 540 °C (700 to 1000 °F)

From 21 to 1090 °C (70 to 2000 °F)

mm/mm/°C × 10−6

in./in./°F × 10−6

mm/mm/°C × 10−6

in./in./°F × 10−6

100 °C (212 °F)

12.8

7.1

...

...

26.0

HA

540 °C 1090 °C (1000 °F) (2000 °F) 27.2

...

HC

11.3

6.3

13.9

7.7

21.8

31.0

41.9

HD

13.9

7.7

16.6

9.2

21.8

31.0

41.9

HE

17.3

9.6

20.0

11.1

14.7

21.5

31.5

HF

17.8

9.9

19.3

10.7

14.4

21.3

...

HH (type I)(a)

17.1

9.5

19.3

10.7

14.2

20.8

30.3

HH (type II)(a)

17.1

9.5

19.3

10.7

14.2

20.8

30.3

HI

17.8

9.9

19.4

10.8

14.2

20.8

30.3

HK

16.9

9.4

18.7

10.4

13.7

20.4

32.2

HL

16.6

9.2

18.2

10.1

14.2

21.1

33.4

HN

16.7

9.3

18.4

10.2

13.0

19.0

29.4

HP

16.6

9.2

19.1

10.6

13.0

19.0

29.4

HT

15.8

8.8

18.0

10.0

12.1

18.7

28.2

HU

15.8

8.8

17.5

9.7

12.1

18.7

28.2

HW

14.2

7.9

HX 14.0 7.8 (a) Type I and II specified per ASTM A 447. Source: Ref 10

16.7

9.3

12.5

19.2

29.4

17.1

9.5

12.5

19.2

29.4

Resistance to Hot-Gas Corrosion. The atmospheres most commonly encountered by heat-resistant cast steel are air, flue gases, and process gases; such gases may be either oxidizing or reducing and may be sulfidizing or carburizing if sulfur or carbon is present. The corrosion of heat-resistant alloys by the environment at elevated temperatures varies significantly with alloy type, temperature, velocity, and the nature of the specific environment to which the part is exposed. Table 14 presents a general ranking of the standard cast, heat-resistant grades in various environments at 980 °C (1800 °F). Corrosion rates at other temperatures are given in Table 12 . Table 14 Corrosion resistance of heat-resistant cast steels at 980 °C (1800 °F) in 100 h tests in various atmospheres Corrosion rating(a) in indicated atmosphere

Air

Oxidizing flue gas(b)

Reducing flue gas(b)

Reducing flue gas(c)

Reducing flue gas (constant temperature)(d)

Reducing flue gas cooled to 150 °C (300 °F) every 12 h(d)

HA

U

U

U

U

U

U

HC

G

G

G

S

G

G

HD

G

G

G

S

G

G

Alloy


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01 Sep 2005

HE

G

G

G

...

G

...

HF

S

G

S

U

S

S

HH

G

G

G

S

G

G

HI

G

G

G

S

G

G

HK

G

G

G

U

G

G

HL

G

G

G

S

G

G

HN

G

G

G

U

S

S

HP

G

G

G

G

G

...

HT

G

G

G

U

S

U

HU

G

G

G

U

S

U

HW

G

G

G

U

U

U

HX G G G S G U (a) G, good (corrosion rate r < 1.27 mm/yr, or 50 mils/yr); S, satisfactory (r < 2.54 mm/yr, or 100 mils/yr); U, unsatisfactory (r > 2.54 mm/yr, or 100 mils/yr). (b) Contained 2 g of sulfur/m3 (5 grains S/100 ft3). (c) Contained 120 g S/m3 (300 grains S/100 ft3). (d) Contained 40 g S/m3 (100 grains S/100 ft3)

Manufacturing Characteristics Foundry practices for cast high-alloy steels for corrosion resistance or heat resistance are essentially the same as those used for cast plain carbon steels. Details on melting practice, metal treatment, and foundry practices, including gating, risering, and cleaning of castings, are available in Casting, Volume 15 of ASM Handbook (formerly 9th Edition Metals Handbook). Iron-base alloys can be cast from heats melted in electric arc furnaces that have either acid or basic linings. When melting is done in acid-lined furnaces, however, chromium losses are high and silicon content is difficult to control, and thus acid-lined furnaces are seldom used. Alloys that contain appreciable amounts of aluminum, titanium, or other reactive metals are melted by induction or electron beam processes under vacuum or a protective atmosphere prior to casting. Welding. As the alloy content of steel castings is increased to produce a fully austenitic structure, welding without cracking becomes more difficult. The fully austenitic low-carbon grades tend to form microfissures adjacent to the weld. This tendency toward microfissuring increases as nickel and silicon contents increase and carbon content decreases. Microfissuring is most evident in coarse-grain alloys with a carbon content of approximately 0.10 to 0.20% and a nickel content exceeding 13%. The microfissuring is reduced by an extremely low sulfur content. In welding these grades, low interpass temperatures, low heat inputs, and peening of the weld to relieve mechanical stresses are all effective. If strength is not a great factor, an initial weld deposit or "buttering of the weld" is also occasionally used. Welding of corrosion-resistant steel castings can be done by shielded metal arc welding, gas tungsten arc welding, gas metal arc welding, and electroslag (submerged arc) welding. Austenitic castings are normally welded without preheat, and are solution annealed after welding. Martensitic castings require preheating to avoid cracking during welding and are given an appropriate postweld heat treatment. Specific conditions for welding specific alloys are listed in Table 15 . When welds are properly made, tensile and yield strengths of the welded joint are similar to those of the unwelded castings (Table 8 ). Elongation is generally lower for specimens taken perpendicular to the weld bead. Table 15 Welding conditions for corrosion-resistant steel castings Preheat ACI designation

Type of electrodes used(a)

°C

°F

Postweld heat treatment

CA-6NM

Same composition

100−150

212−300

590−620 °C (1100−1150 °F)

CA-15

410

200−315

400−600

610−760 °C (1125−1400 °F), air cool

CA-40

410 or 420

200−315

400−600

610−760 °C (1125−1400 °F), air cool

CB-7Cu

Same composition or 308

CB-30

442

315−425

600−800

790 °C (1450 °F) min, air cool

CC-50

446

200−700

400−1300

900 °C (1650 °F), air cool

CD-4MCu

Same composition

Not required

Heat to 1120 °C (2050 °F), cool to 1040 °C (1900 °F), quench

CE-30

312

Not required

Quench from 1090−1120 °C (2000−2050 °F)

CF-3

308L

Not required

Usually unnecessary

CF-8

308

Not required

Quench from 1040−1120 °C (1900−2050 °F)

CF-8C

347

Not required

Usually unnecessary


Not required

480−590 °C (900−1100 °F), air cool

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01 Sep 2005

CF-3M

316L

Not required

Usually unnecessary

CF-8M

316

Not required

Quench from 1070−1150 °C (1950−2100 °F)

CF-12M

316

Not required

Quench from 1070−1150 °C (1950−2100 °F)

CF-16F

308 or 308L

Not required

Quench from 1090−1150 °C (2000−2100 °F)

CF-20

308

Not required

Quench from 1090−1150 °C (2000−2100 °F)

CG-8M

317

Not required

Quench from 1040−1120 °C (1900−2050 °F)

CH-20

309

Not required

Quench from 1090−1150 °C (2000−2100 °F)

CK-20

310

CN-7M

320

Quench from 1090−1180 °C (2000−2150 °F)

Not required 200

400

Quench from 1120 °C (2050 °F)

Note: Metal arc, inert-gas arc, and electroslag welding methods can be used. Suggested electrical settings and electrode sizes for various section thicknesses are: Section thickness, mm (in.)

Electrode diameter, mm (in.)

Current, A

Maximum arc voltage, V

2.4 (3=32 )

45−70

24

3.2−6.4 (1=8−1=4) 3.2−6.4 (1=8−1=4) 3.2−6.4 ( =8− =4) 1

1

6.4−13 ( =4− =2) 1

1

≥13 ( =2) (a) Lime-coated electrodes are recommended. 1

1

3.2 ( =8)

70−105

25

5

100−140

25

3

130−180

26

210−290

27

4.0 ( =32 ) 4.8 ( =16 ) 1

6.4 ( =4)

Most of the corrosion-resistant cast steels, such as the CF-8 or CF-8M grade, are readily weldable, especially if their microstructures contain small percentages of δ-ferrite. Because stainless steels can become sensitized and lose their corrosion resistance if subjected to temperatures above 425 °C (800 °F), great care must be taken in welding to make certain that the casting or fabricated component is not heated excessively. For this reason, many stainless steels are almost never preheated. In many cases, the weld is cooled with a water spray between passes to reduce the interpass temperature to 150 °C (300 °F) or below. Any welding performed on the corrosion-resistant grades will affect the corrosion resistance of the casting, but for many services the castings will perform satisfactorily in the as-welded condition. Where extremely corrosive conditions exist or where SCC may be a problem, complete reheat treatment may be required after welding. Heating the casting above 1065 °C (1950 °F) and then cooling it rapidly redissolves the carbides precipitated during the welding operation and restores corrosion resistance. When maximum corrosion resistance is desired and postweld heat treatment (solution annealing) cannot be performed, alloying elements can be added to form stable carbides. Although niobium and titanium both form stable carbides, titanium is readily oxidized during the casting operation and therefore is seldom used. The niobium-stabilized grade CF-8C is the most commonly used cast grade. The stability of the niobium carbides prevents the formation of chromium carbides and the consequent chromium depletion of the base metal. This grade may therefore be welded without postweld heat treatment. Another approach to take when postweld heat treatment is undesirable or impossible is to keep the carbon content below 0.03%, as in the CF-3 and CF-3M grades. At this low carbon level, the depletion of the chromium due to carbide precipitation is so slight that the corrosion resistance of the grade is unaffected by the welding operation.

Galling Stainless steel castings are susceptible to galling and seizing when dry surfaces slide or chafe against each other. However, the surfaces of the castings can be nitrided so that they are hard and wear resistant. Tensile properties are not impaired. Nitriding reduces resistance to corrosion by concentrated nitric or mixed acids. Parts such as gate disks for gate valves are usually furnished in the solution-treated condition, but may be nitrided to reduce susceptibility to seizing in service. Similar results are obtained by hardfacing with cobalt-chromium-tungsten alloys.

Magnetic Properties The magnetic properties of high-alloy castings depend on microstructure. The straight chromium types are ferritic and ferromagnetic. All other grades are mainly austenitic, with or without minor amounts of ferrite, and are either weakly magnetic or wholly nonmagnetic. Cast nonmagnetic parts for applications in radar and in minesweepers require close control of ferrite content. Thicker sections have higher permeability than thinner sections. Therefore, to ensure low magnetic permeability in all areas of a casting, magnetic permeability checks should be made on the thicker sections. REFERENCES 1. M. Prager, Cast High Alloy Metallurgy, in Steel Casting Metallurgy, J. Svoboda, Ed., Steel Founders' Society of America,


Page 1441



01 Sep 2005

1984, p 221-245 2. C.E. Bates and L.T. Tillery, Atlas of Cast Corrosion-Resistant Alloy Microstructures, Steel Founders' Society of America, 1985 3. F. Beck, E.A. Schoefer, E. Flowers, and M. Fontana, New Cast High Strength Alloy Grades by Structure Control, in Advances in the Technology of Stainless Steels and Related Alloys, STP 369, American Society for Testing and Materials, 1965, p 159-174 4. T.M. Devine, Mechanism of Intergranular Corrosion and Pitting Corrosion of Austenitic and Duplex 308 Stainless Steel, J. Electrochem. Soc., Vol 126 (No. 3), 1979, p 374 5. E.E. Stansbury, C.D. Lundin, and S.J. Pawel, Sensitization Behavior of Cast Stainless Steels Subjected to Simulated Weld Repair, in Proceedings of the 38th SFSA Technical and Operating Conference, Steel Founders' Society of America, 1983, p 223 6. S. Shimodaira et al., Mechanisms of Transgranular Stress Corrosion Cracking of Duplex and Ferrite Stainless Steels, in Stress Corrosion Cracking and Hydrogen Embrittlement in Iron Base Alloys, NACE Reference Book 5, National Association of Corrosion Engineers, 1977 7. P.L. Andersen and D.J. Duquette, The Effect of Cl− Concentration and Applied Potential on the SCC Behavior of Type 304 Stainless Steel in Deaerated High Temperature Water, Corrosion, Vol 36 (No. 2), 1980, p 85−93 8. S.B. Shendye, "Effect of Long Term Elevated Temperature Exposure on the Mechanical Properties and Weldability of Cast Duplex Steels," Master's thesis, Oregon Graduate Center, 1985 9. High Alloy Data Sheet, Heat Series, in Steel Castings Handbook Supplement 9, Steel Founders' Society of America 10. "Heat and Corrosion-Resistant Castings," The International Nickel Company, 1978 11. A. Brasunas, J.T. Glow, and O.E. Harder, Resistance of Fe-Ni-Cr Alloys to Corrosion in Air at 1600 to 2200 °F, in Proceedings of the ASTM Symposium for Gas Turbines, American Society for Testing and Materials, 1946, p 129−152


Page 1442


Elevated-Temperature Properties of Stain...

01 Sep 2005

ASM Handbook, Volume 1, Properties and Selection: Irons, Steels, and High Performance Alloys Section: Specialty Steels and Heat-Resistant Alloys Elevated-Temperature Properties of Stainless Steels STAINLESS STEELS are widely used at elevated temperatures when carbon and low-alloy steels do not provide adequate corrosion resistance and/or sufficient strength at these temperatures. Carbon and low-alloy steels are generally more economical than stainless steels and are often used in applications with temperatures below about 370 °C (700 °F). Several low-alloy steels with moderate chromium contents (between 1 and 10%) and improved high-temperature strength are also widely used at elevated temperatures above 370 °C (700 °F). These steels include the creep-resistant chromium-molybdenum ferritic steels discussed in the article"Elevated-Temperature Properties of Ferritic Steels" in this Volume. Carbon steels may even be suitable for temperatures above 370 °C (700 °F), if high strength and oxidation are not concerns. This article deals with the wrought stainless steels used for high-temperature applications (see the article "Cast Stainless Steels" in this Volume for the elevated-temperature properties of cast stainless steels). Corrosion resistance is often the first criterion used to select stainless steel for a particular application. However, strength is also a significant factor in a majority of elevated-temperature applications and may even be the key factor governing the choice of a stainless steel. The stainless steels used in applications in which high-temperature strength is important are sometimes referred to as heat-resistant steels. Table 1 gives some typical compositions of wrought heat-resistant stainless steels, which are grouped into ferritic, martensitic, austenitic, and precipitation-hardening grades. Of these steels, the austenitic grades offer the highest strength at high temperatures (Fig. 1 ). The precipitation-hardening steels have the highest strength at lower temperatures (Fig. 1 ), but they weaken considerably at temperatures above about 425 °C (800 °F). Table 1 Nominal compositions of wrought iron-base heat-resistant alloys Designation

UNS number

Composition, % C

Cr

Ni

Mo

N

Nb

Ti

Other

Ferritic stainless steels 405

S40500

0.15 max

13.0

...

...

...

...

...

0.2 Al

406

...

0.15 max

13.0

...

...

...

...

...

4.0 Al

409

S40900

0.08 max

11.0

0.5 max

...

...

...

6 × C min

...

429

S42900

0.12 max

15

...

...

...

...

...

...

430

S43000

0.12 max

16.0

...

...

...

...

...

...

434

S43400

0.12 max

17.0

...

1.0

...

...

...

...

439

S43035

0.07 max

18.25

...

...

...

...

12 × C min

18 SR

1.10 Ti max

...

0.05

18.0

...

...

...

...

0.40 max

18Cr-2Mo

S44400

...

18.5

...

2.0

...

(a)

(a)

2.0 Al max

446

S44600

0.20 max

25.0

...

...

0.25

...

...

...

E-Brite 26-1

S44627

0.01 max

26.0

...

1.0

0.015 max

0.1

...

...

26-1Ti

S44626

0.04

26.0

...

1.0

...

...

10 × C min

...

29Cr-4Mo

S44700

0.01 max

29.0

...

4.0

0.02 max

...

...

...

0.8 (Ti + Nb) max

Quenched and tempered martensitic stainless steels


Page 1443



01 Sep 2005

403

S40300

0.15 max

12.0

...

...

...

...

...

...

410

S41000

0.15 max

12.5

...

...

...

...

...

...

410Cb

S41040

0.15 max

12.5

...

...

...

0.12

...

...

416

S41600

0.15 max

13.0

...

0.6(b)

...

...

...

0.15 min S

422

S42200

0.20

12.5

0.75

1.0

...

...

...

1.0 W, 0.22 V

H-46

...

0.12

10.75

0.50

0.85

0.07

0.30

...

0.20 V

Moly Ascoloy

...

0.14

12.0

2.4

1.80

0.05

...

...

0.35 V

Greek Ascoloy

S41800

0.15

13.0

2.0

...

...

...

...

3.0 W

Jethete M-152

...

0.12

12.0

2.5

1.7

...

...

...

0.30 V

Almar 363

...

0.05

11.5

4.5

...

...

...

10 × C min

...

431

S43100

0.20 max

16.0

2.0

...

...

...

...

...

Lapelloy

S42300

0.30

11.5

...

2.75

...

...

...

0.25 V

0.75

...

8 × C min

...

1.5 Cu

Precipitation-hardening martensitic stainless steels Custom 450 Custom 455

...

0.05 max

15.5

6.0

...

0.03

11.75

8.5

...

...

0.30

1.2

2.25 Cu

15-5 PH

S15500

0.07

15.0

4.5

...

...

0.30

...

3.5 Cu

17-4PH

S17400

0.04

16.5

4.25

...

...

0.25

...

3.6 Cu

PH 13-8 Mo

S13800

0.05

12.5

8.0

2.25

...

...

...

1.1 Al

Precipitation-hardening semiaustenitic stainless steels AM-350

S35000

0.10

16.5

4.25

2.75

0.10

...

...

...

AM-355

S35500

0.13

15.5

4.25

2.75

0.10

...

...

...

17-7 PH

S17700

0.07

17.0

7.0

...

...

...

...

1.15 Al

PH 15-7 Mo

S15700

0.07

15.0

7.0

2.25

...

...

...

1.15 Al

Austenitic stainless steels 304

S30400

0.08 max

19.0

10.0

...

...

...

...

...

304H

S30409

0.04−0. 10

19.0

10.0

...

...

...

...

...

304L

S30403

0.03 max

19.0

10.0

...

...

...

...

...

304N

S30451

0.08 max

19.0

9.25

...

0.13

...

...

...

309

S30900

0.30 max

23.0

13.0

...

...

...

...

...

309H

S30909

0.04−0. 10

23.0

13.0

...

...

...

...

...

310

S31000

0.25 max

25.0

20.0

...

...

...

...

...

310H

S31009

0.04−0. 10

25.0

20.0

...

...

...

...

...

316

S31600

0.08 max

17.0

12.0

2.5

...

...

...

...

316L

S31603

0.03 max

17.0

12.0

2.5

...

...

...

...

316N

S31651

0.08 max

17.0

12.0

2.5

0.13

...

...

...

316H

S31609

0.04−0. 10

17.0

12.0

2.5

...

...

...

...


Page 1444



01 Sep 2005

316LN

S31653

0.035 max

17.0

12.0

2.5

0.13

...

...

...

317

S31700

0.08 max

19.0

13.0

3.5

...

...

...

...

317L

S31703

0.035 max

19.0

13.0

3.5

...

...

...

...

321

S32100

0.08 max

18.0

10.0

...

...

...

5 × C min, 0.70 max

...

321H

S32109

0.04−0. 10

18.0

10.0

...

...

...

4 × C min, 0.60 max

...

347

S34700

0.08 max

18.0

11.0

...

...

10 × C min(c)

...

1.0 (Nb + Ta) max

347H

S34709

0.04−0. 10

18.0

11.0

...

...

8 × C min(c)

...

1.0 (Nb + Ta) max

348

S34800

0.08 max

18.0

11.0

...

...

10 × C min(c)

...

0.10 Ta max, 1.0 (Nb + Ta) max

348H

S34809

0.04−0. 10

18.0

11.0

...

...

8 × C min(c)

...

0.10 Ta max, 1.0 (Nb + Ta) max

19-9 DL

K63198

0.30

19.0

9.0

1.25

...

0.4

0.3

1.25 W

19-9 DX

K63199

0.30

19.2

9.0

1.5

...

...

0.55

1.2 W

...

0.12

16.0

14.0

2.5

...

0.4

0.3

3.0 Cu

201

S20100

0.15 max

17

4.2

...

0.25 max

...

...

202

S20200

0.09

18.0

5.0

...

0.10

...

...

205

S20500

0.18

17.2

1.4

...

0.36

...

...

17-14-CuMo

... 8.0 Mn ...

216

S21600

0.05

20.0

6.0

2.5

0.35

...

...

8.5 Mn

21-6-9

S21900

0.04 max

20.25

6.5

...

0.30

...

...

9.0 Mn

Nitronic 32

S24100

0.10

18.0

1.6

...

0.34

...

...

12.0 Mn

Nitronic 33

S24000

0.08 max

18.0

3.0

...

0.30

...

...

13.0 Mn

Nitronic 50

...

0.06 max

21.0

12.0

2.0

0.30

0.20

...

5.0 Mn

Nitronic 60

S21800

0.10 max

17.0

8.5

2.0

...

...

...

8.0 Mn, 0.20 V, 4.0 Si

...

...

16.0 Mn, 0.40 Si, 1.0 Cu

Carpenter 18-18 S28200 0.10 18.0 50%) because γ′ shearing is the primary strengthening mechanism. With the mean free edge-to-edge distance in the γ matrix between the precipitates being smaller than the average precipitate size itself, dislocation shearing of the γ′ particle is favored over Orowan dislocation looping around the γ′ particles. Fig. 12 Influence of alloying elements on the lattice parameter of binary nickel alloys. Source: Ref 21


Page 1559


Directionally Solidified and Single-Crys...

01 Sep 2005

Detailed transmission electron microscopy studies of dislocation movement in cast high-strength superalloys, such as MAR-M 002 (Table 2 ) and its single-crystal derivative SRR 99 (Table 1 ), have shown the importance of ensuring that the antiphase boundary (APB) energy is high, so that the stacking fault mode of creep deformation occurs at temperatures up to 850 °C (1562 °F), thus ensuring high creep strength (Ref 17). Tantalum additions raise the APB energy relative to the stacking fault energy (Ref 17), leading to the increased tendency for stacking faults to be formed at lower temperatures. The CMSX-2 alloy is designed to provide good SX foundry performance because castability is a crucial alloy performance criterion for any complex, thin-wall turbine blade or vane component, a characteristic sometimes given limited attention in alloy design. It affects not only the yield and cost of components but also the defect level and therefore component performance. Single-crystal casting defects of concern are: • • • • •

Freckling: A spiral of equiaxed grains caused by elemental segregation in the liquid state Slivers: Moderate-angle grain defects Microporosity: A uniform distribution of interdendritic micropores Spurious grains: High-angle grain boundaries Stable oxide inclusions: Al2O3

• Carbides: TiC The partial substitution of tantalum for tungsten in the CMSX-2 alloy, compared to the MAR-M 247 chemistry, helps overcome the freckling problems inherent in the low-tantalum, high-tungsten single-crystal alloys. The strong γ′-forming elements, aluminum and titanium, which are also low density, tend to segregate to the last liquid to solidify in the interdendritic spaces during the SX solidification process. This can create density changes and consequential flow in the liquid metal close to the solidification front, which can nucleate freckle trails of equiaxed grains. This can occur particularly under conditions of low


Page 1560



01 Sep 2005

or changing thermal gradients. Tantalum, which is a strong γ′-forming element of high density, also tends to segregate to the last liquid to solidify in the interdendritic spaces and thus evens out these density changes in the liquid, or mushy, zone and reduces freckling tendencies. Several studies undertaken in the United States, Europe, and Japan confirm that high [N] and [O] levels in single-crystal superalloy ingot adversely affect SX casting grain yield, supporting the importance for low [N] and [O] levels in the master alloy. Carbon, sulfur, and [O] master alloy impurities are shown to transfer nonmetallic inclusions, such as Al2O3, (Ti,Ta) C/N, and (Ti,Ta)x S, to SX parts (Ref 22). Grain defects can nucleate on these inclusions. Several second-generation, rhenium-containing, single-crystal superalloys have been developed for turbine engine applications. Two typical compositions are given in Table 3 . Rhenium partitions mainly to the γ matrix; this retards coarsening of the γ′-strengthening phase and increases γ/γ′ misfit (Ref 23). Atom-probe microanalysis of rhenium-containing modifications of the PWA 1480 and CMSX-2 alloys reveals the occurrence of short-range order in the matrix with small rhenium clusters (~1.0 nm, or 10 Å, in size) detected in the γ in the alloys (Ref 24). The rhenium clusters can act as more efficient obstacles against dislocation movement compared to isolated solute atoms in the γ solid solution; therefore, they play a significant role in improving the creep strength. The Larson-Miller stress-rupture comparison of CMSX-4 and CMSX-2/3 is shown in Fig. 13 . The stress-rupture temperature capability advantage of CMSX-4 over CMSX-2/3 is 27 °C (48 °F) (density corrected) in the 248 MPa/982 °C (36 ksi/1800 °F) region. In the 103 MPa/1121 °C (15 ksi/2050 °F) region, the stress-rupture temperature capability advantage is 30 °C (54 °F) (density corrected). The data also indicate that CMSX-4 has a potential peak-use temperature under stress of at least 1149 °C (2100 °F). Fig. 13 Larson-Miller stress-rupture strength of CMSX-4 versus CMSX-2/3

Single-Crystal Casting Techniques. A variety of single-crystal airfoil component-casting techniques have been developed to production status around the world in the last 10 years. Most involve a withdrawal-type vacuum induction casting furnace with mold susceptor heating. Cooling plate sizes range in diameter from 140 to 610 mm (51=2to 24 in.). Some of the developed SX casting techniques are presented in Ref 12, 13, 25, and 26. The modern helicopter engine turbine vane shown in Fig. 14 represents a difficult cored configuration. The large shrouds and core make this vane susceptible to shrinkage, grain nucleation, and recrystallization during solution heat treatment. Single-crystal casting processes developed by the Allison Gas Turbine Division of General Motors Corporation result in high yields for this vane in CMSX-3. Similar yields have been demonstrated with CMSX-4 using the same Allison casting process. Fig. 14 SX turbine vane cast in CMSX-4


Page 1561



01 Sep 2005

Single-Crystal Heat Treatment and Microstructures. With regard to solutioning, the latest multistep ramped cycles developed for single-crystal components are designed to completely solution the γ′ and most of the γ/γ′ eutectic without incipient melting. An additional benefit of the high-temperature cycles is the element homogenization effect, as shown in Fig. 15 . Alloy CMSX-4, which is solutioned at a maximum temperature of 1321 °C (2410 °F) in commercial vacuum heat treatment furnaces, readily attains the 99%+ ( greater than ¿ much greater than ≥ greater than or equal to ∞ infinity


Page 1616


Abbreviations, Symbols, and Tradenames

01 Sep 2005

∝ is proportional to; varies as ∫ integral of < less than ° much less than ≤ less than or equal to ± maximum deviation − minus; negative ion charge × diameters (magnification); multiplied by · multiplied by / per % percent + plus; positive ion charge √ square root of ~ approximately; similar to ∂ partial derivative ∆ change in quantity; an increment; a range ε strain ²_ strain rate µ friction coefficient; magnetic permeability µin. microinch µm micron (micrometer) µs microsecond ν Poisson's ratio π pi (3.141592) ρ density σ stress Σ summation of τ shear stress Ω ohm Greek Alphabet Α, α alpha Β, β beta Γ, γ gamma ∆ δ delta Ε, ε epsilon Ζ, ζ zeta Η, η eta Θ, θ theta Ι, ι iota Κ, κ kappa Λ, λ lambda Μ, µ mu Ν, ν nu Ξ, ξ xi Ο, ο omicron Π, π pi Ρ, ρ rho Σ, σ sigma Τ, τ tau Υ, υ upsilon Χ, χ chi Ψ, ψ psi Ω, ω omega Tradenames


Page 1617


Abbreviations, Symbols, and Tradenames

01 Sep 2005

AF-56 is a registered tradename of Allison Gas Turbine, Division of General Motors Corporation. AL-6X, AL-6XN, AL 29-4C, AL 29-4-2, AL 904L, AL 2205, ALFA IV, E-Brite, 26-1, Sealmet, and 203 EZ are registered tradenames of Allegheny Ludlum Steel, Division of Allegheny Ludlum Corporation. AM1 is a registered tradename of SNECMA/ONERA. CM 247 LC and CMSX are registered tradenames of Cannon-Muskegon Corporation. Cronifer is a registered tradename of Vereingte Deutsche Metallwerks. Cryogenic Tenelon and Tenelon are registered tradenames of USS, Division of USX Corporation. Custom 450, Custom 455, Gall-Tough, Pyromet, TrimRite, 7-Mo PLUS, 18-18 PLUS, 20Cb-3, 20Mo-4 and 20Mo-6 are registered tradenames of Carpenter Technology Corporation. Discaloy is a registered tradename of Westinghouse Electric Corporation. DP3 is a registered tradename of Sumitomo Metal America, Inc. Esshete is a registered tradename of British Steel Corporation. Ferralium is a registered tradename of Bonar Langley Alloy Ltd. Hastelloy and Haynes are registered tradenames of Haynes International, Inc. Incoloy, Inconel, Nimocast, and Nimonic are registered tradenames of INCO Alloys International, Inc. JS700 is a registered tradename of Jessop Steel Company. MAR-M is a registered tradename of Martin Marietta Corporation. Monit is a registered tradename of Uddeholms Aktiebolag. MP (Multiphase) is a registered tradename of Standard Pressed Steel Company. Nitronic and PH 13-8 Mo are registered tradenames of Baltimore Specialty Steels Corporation. PH 15-7 MO, 12SR, 15-5 PH, 17-4 PH, 18 SR, and 21-6-9 are registered tradenames of Armco Advanced Materials Corporation. PWA 1484 is a registered tradename of Pratt & Whitney Aircraft. RA85H is a registered tradename of Rolled Alloys, Inc. René is a registered tradename of General Electric Company. René 41 is a registered tradename of Allvac Metals Company, a Teledyne Company. RR 2000 and SRR 99 are registered tradenames of Rolls Royce, Inc. Sanicro and 3RE60 are registered tradenames of Sandvik, Inc. Sea-Cure is a registered tradename of Crucible, Inc. Stellite is a registered tradename of Deloro Stellite, Inc. Udimet is a registered tradename of Special Metals Corporation. Unitemp is a registered tradename of Universal Cyclops Steel Corporation. Uranus is a registered tradename of Compagnie des Ateliers et Forges de la Loire. Vitallium is a registered tradename of Pfizer Hospital Products Group, Inc. Waspaloy is a registered tradename of United Technologies, Inc. 253MA and 254SMO are registered tradenames of Avesta Stainless, Inc.


Page 1618

Properties and Selection: Irons, Steels, and High Performance Alloys

Metallographer's Guide: Practices and Procedures for Irons and Steels

Steels Processing Structure and Performance

Steels: Processing, Structure, And Performance

Heat Treater's Guide: Practices and Procedures for Irons and Steels

Microstructure of Steels and Cast Irons (Engineering Materials and Processes)

Duplex Stainless Steels, Microstructure, properties and applications

Steels: Microstructure and Properties, Third Edition

Alloys: Preparation, Properties, Applications

Structure and properties of engineering alloys

Structure and Properties of Nonferrous Alloys

High Performance Organic Coatings: Selection, Application and Evaluation

Engineering Properties and Applications of Lead Alloys

Nanocomposites: Preparation, Properties and Performance

High Temperature Strain of Metals and Alloys

High Temperature Strain Of Metals And Alloys

High Performance Instrumentation and Automation

Engineering and High Performance Plastics

Nanostructured metals and alloys - Processing, microstructure, mechanical properties and applications

The Alloy Tree: A Guide to Low-alloy Steels, Stainless Steels and Nickel-base Alloys

Columns for Gas Chromatography: Performance and Selection

Steels: metallurgy and applications

The Alloy Tree: A Guide to Low-alloy Steels, Stainless Steels and Nickel-base Alloys

Maraging Steels: Modelling of Microstructure, Properties and Applications

Stainless Steels. Their properties and their suitability for welding

Running Irons

Powder Metallurgy Stainless Steels: Processing, Microstructures, and Properties

Microstructure and Properties of High-Temperature SuperConductors

Running Irons

Properties and Selection: Irons, Steels, and High Performance Alloys

Metallographer's Guide: Practices and Procedures for Irons and Steels

Steels Processing Structure and Performance

Steels: Processing, Structure, And Performance

Heat Treater's Guide: Practices and Procedures for Irons and Steels

Microstructure of Steels and Cast Irons (Engineering Materials and Processes)

Duplex Stainless Steels, Microstructure, properties and applications

Steels: Microstructure and Properties, Third Edition

Alloys: Preparation, Properties, Applications

Structure and properties of engineering alloys

Structure and Properties of Nonferrous Alloys

High Performance Organic Coatings: Selection, Application and Evaluation

Engineering Properties and Applications of Lead Alloys

Nanocomposites: Preparation, Properties and Performance

High Temperature Strain of Metals and Alloys

High Temperature Strain Of Metals And Alloys

High Performance Instrumentation and Automation

Engineering and High Performance Plastics

Nanostructured metals and alloys - Processing, microstructure, mechanical properties and applications

The Alloy Tree: A Guide to Low-alloy Steels, Stainless Steels and Nickel-base Alloys

Columns for Gas Chromatography: Performance and Selection

Steels: metallurgy and applications

The Alloy Tree: A Guide to Low-alloy Steels, Stainless Steels and Nickel-base Alloys

Maraging Steels: Modelling of Microstructure, Properties and Applications

Stainless Steels. Their properties and their suitability for welding

Running Irons

Powder Metallurgy Stainless Steels: Processing, Microstructures, and Properties

Microstructure and Properties of High-Temperature SuperConductors

Running Irons

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