Uhlig's Corrosion Handbook, 3rd Edition (The ECS Series of Texts and Monographs)

UHLIG’S CORROSION HANDBOOK THE ELECTROCHEMICAL SOCIETY SERIES ECS-The Electrochemical Society 65 South Main Street P...

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UHLIG’S CORROSION HANDBOOK

THE ELECTROCHEMICAL SOCIETY SERIES

ECS-The Electrochemical Society 65 South Main Street Pennington, NJ 08534-2839 http://www.electrochem.org

Corrosion Handbook Edited by Herbert H. Uhlig Modern Electroplating, Third Edition Edited by Frederick A. Lowenheim Modern Electroplating, Fifth Edition Edited by Mordechay Schlesinger and Milan Paunovic The Electron Microprobe Edited by T. D. McKinley, K. F. J. Heinrich, and D. B. Wittry Chemical Physics of Ionic Solutions Edited by B. E. Conway and R. G. Barradas High-Temperature Materials and Technology Edited by Ivor E. Campbell and Edwin M. Sherwood Alkaline Storage Batteries S. Uno Falk and Alvin J. Salkind The Primary Battery (in Two Volumes) Volume I Edited by George W. Heise and N. Corey Cahoon Volume II Edited by N. Corey Cahoon and George W. Heise Zinc-Silver Oxide Batteries Edited by Arthur Fleischer and J. J. Lander Lead-Acid Batteries Hans Bode Translated by R. J. Brodd and Karl V. Kordesch Thin Films-Interdiffusion and Reactions Edited by J. M. Poate, M. N. Tu, and J. W. Mayer Lithium Battery Technology Edited by H. V. Venkatasetty Quality and Reliability Methods for Primary Batteries P. Bro and S. C. Levy

Techniques for Characterization of Electrodes and Electrochemical Processes Edited by Ravi Varma and J. R. Selman Electrochemical Oxygen Technology Kim Kinoshita Synthetic Diamond: Emerging CVD Science and Technology Edited by Karl E. Spear and John P. Dismukes Corrosion of Stainless Steels A. John Sedriks Semiconductor Wafer Bonding: Science and Technology Q.-Y. Tong and U. G€oscle Uhlig’s Corrosion Handbook, Second Edition Edited by R. Winston Revie Atmospheric Corrosion Christofer Leygraf and Thomas Graedel Electrochemical Systems, Third Edition John Newman and Karen E. Thomas-Alyea Fundamentals of Electrochemistry, Second Edition V. S. Bagotsky Fundamentals of Electrochemical Deposition, Second Edition Milan Paunovic and Mordechay Schlesinger Electrochemical Impedance Spectroscopy Mark E. Orazem and Bernard Tribollet Fuel Cells: Problems and Solutions Vladimir S. Bagotsky

UHLIG’S CORROSION HANDBOOK THIRD EDITION

Edited by

R. WINSTON REVIE CANMET Materials Technology Laboratory Ottawa, Ontario, Canada

Copyright Ó 2011 by John Wiley & Sons, Inc. All rights reserved Published by John Wiley & Sons, Inc., Hoboken, New Jersey Published simultaneously in Canada No part of this publication may be reproduced, stored in a retrieval system, or transmitted in any form or by any means, electronic, mechanical, photocopying, recording, scanning, or otherwise, except as permitted under Section 107 or 108 of the 1976 United States Copyright Act, without either the prior written permission of the Publisher, or authorization through payment of the appropriate per-copy fee to the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, (978) 750-8400, fax (978) 750-4470, or on the web at www.copyright.com. Requests to the Publisher for permission should be addressed to the Permissions Department, John Wiley & Sons, Inc., 111 River Street, Hoboken, NJ 07030, (201) 748-6011, fax (201) 748-6008, or online at http://www.wiley.com/go/permission. Limit of Liability/Disclaimer of Warranty: While the publisher and author have used their best efforts in preparing this book, they make no representations or warranties with respect to the accuracy or completeness of the contents of this book and specifically disclaim any implied warranties of merchantability or fitness for a particular purpose. No warranty may be created or extended by sales representatives or written sales materials. The advice and strategies contained herein may not be suitable for your situation. You should consult with a professional where appropriate. Neither the publisher nor author shall be liable for any loss of profit or any other commercial damages, including but not limited to special, incidental, consequential, or other damages. For general information on our other products and services or for technical support, please contact our Customer Care Department within the United States at (800) 762-2974, outside the United States at (317) 572-3993 or fax (317) 572-4002. Wiley also publishes its books in a variety of electronic formats. Some content that appears in print may not be available in electronic formats. For more information about Wiley products, visit our web site at www.wiley.com. Library of Congress Cataloging-in-Publication Data: Uhlig’s corrosion handbook / edited by R. Winston Revie–3rd ed. p. cm. –(The ECS series of texts and monographs) Includes index. ISBN 978-0-470-08032-0

Printed in the United States of America eBook ISBN: 978-0-470-87285-7 oBook ISBN: 978-0-470-87286-4 10 9 8

7 6 5

4 3 2 1

Herbert H. Uhlig March 3, 1907–July 3, 1993 This Handbook is dedicated to the memory of Herbert H. Uhlig. Herbert Uhlig began his career at MIT in 1936 where, with the exception of the interruption caused by World War II, he remained until his retirement nearly 40 years later, bringing the MIT Corrosion Laboratory to a level of international prominence that it retains to this day as a major center of excellence. He helped to establish the Corrosion Division of The Electrochemical Society in 1942 and served as President of the Society in 1955–1956. His characteristics as an uncompromising innovator and meticulous scientist who insisted on reliable data and on achieving results led to the success of his many endeavors as educator and mentor, including the Corrosion Handbook, published in 1948, that he conceived, organized, and edited.

EDITORIAL ADVISORY BOARD

Robert Baboian RB Corrosion Service Greenville, Rhode Island

Tomomi Murata Japan Science and Technology Agency Saitama, Japan

Donald L. Jordan Ford Motor Company Dearborn, Michigan

Srdjan Nesic Ohio University Athens, Ohio

Russell D. Kane iCorrosion LLC Houston, Texas

Howard W. Pickering The Pennsylvania State University University Park, Pennsylvania

Martin Kendig Rockwell International Science Center Thousand Oaks, California

Pierre R. Roberge Royal Military College of Canada Kingston, Ontario Canada

Pradip R. Khaladkar E.I. DuPont de Nemours & Co., Inc. Wilmington, Delaware

Guenter A. Schmitt Iserlohn University of Applied Sciences Iserlohn, Germany

Ronald M. Latanision Exponent-Failure Analysis Associates, Inc. Natick, Massachusetts

Roger W. Staehle University of Minnesota and Industrial Consultant North Oaks, Minnesota

Barry MacDougall National Research Council of Canada Ottawa, Ontario, Canada

Graham Wood University of Manchester Manchester, UK

CONTENTS

Foreword

xvii

Foreword to the Second Edition

xix

Preface

xxi

Contributors

xxiii

Introductory Notes on Climate Change, Life-Cycle Design, and Corrosion of Steel T. Murata

xxvii

PART I 1

BASICS OF CORROSION SCIENCE AND ENGINEERING

Corrosion Failure Analysis with Case Histories

3

L. E. Eiselstein and R. Huet

2

Cost of Metallic Corrosion

15

J. Kruger

3

Economics of Corrosion

21

E. D. Verink, Jr.

4

Lifetime Prediction of Materials in Environments

31

R. W. Staehle

5

Estimating the Risk of Pipeline Failure Due to Corrosion

75

M. Nessim

6

Designing to Prevent Corrosion

83

E. D. Verink, Jr.

7

Simplified Procedure for Constructing Pourbaix Diagrams

93

E. D. Verink, Jr. ix

x

CONTENTS

8 Pourbaix Diagrams for Multielement Systems

103

W. T. Thompson, M. H. Kaye, C. W. Bale, and A. D. Pelton

9 Computation of Pourbaix Diagrams at Elevated Temperature

111

M. H. Kaye and W. T. Thompson

10 Galvanic Corrosion

123

X. G. Zhang

11 Dealloying

145

R. Heidersbach

12 Passivity

151

J. Kruger

13 Localized Corrosion of Passive Metals

157

H. Bo¨hni

14 Stress Corrosion Cracking

171

R. N. Parkins

15 Hydrogen-Induced Cracking and Sulfide Stress Cracking

183

M. Elboujdaini

16 Corrosion Fatigue

195

Y.-Z. Wang

17 Flow Effects on Corrosion

203

K. D. Efird

18 Erosion–Corrosion in Single- and Multiphase Flow

215

J. Postlethwaite and S. Nesˇic´

19 Carbon Dioxide Corrosion of Mild Steel

229

S. Nesˇic´

20 High-Temperature Oxidation

247

C. A. C. Sequeira

21 Thermochemical Evaluation of Corrosion Product Stabilities for Alloys in Gases at High Temperature

281

W. T. Thompson, R. C. John, and A. L. Young

22 A Procedure to Compute Equilibrium Gas-Phase Speciation for Use with Predominance Diagrams

287

M. H. Piro, B. J. Lewis, and W. T. Thompson

23 Atmospheric Corrosion P. R. Roberge

299

CONTENTS

24

Atmospheric Corrosion in Cold Regions

327

G. A. King

25

Corrosion by Soils

333

T. R. Jack and M. J. Wilmott

26

Microbial Degradation of Materials: General Processes

351

J.-D. Gu, T. E. Ford, and R. Mitchell

27

Corrosion Probability and Statistical Evaluation of Corrosion Data

365

T. Shibata

PART II 28

NONMETALS

Corrosion of Refractories and Ceramics

387

M. Rigaud

29

Corrosion of Glass

399

B. Grambow

30

Microbiological Degradation of Polymeric Materials

421

J.-D. Gu, T. E. Ford, D. B. Mitton, and R. Mitchell

31

Durability of Concrete

439

V. M. Malhotra

32

Microbiological Corrosion of Concrete

451


33

Microbial Degradation of Wood

461

P. I. Morris

34

Use of Chemicals to Prevent Degradation of Wood

469

J. N. R. Ruddick

PART III 35

METALS

Metal–Matrix Composites

481

L. H. Hihara

36

Environmental Degradation of Engineered Barrier Materials in Nuclear Waste Repositories

503

R. B. Rebak

37

Corrosion Behavior of Electrodeposited Nanocrystals

517

U. Erb

38

Corrosion of Shape Memory and Superelastic Alloys L. E. Eiselstein

529

xi

xii

CONTENTS

39 Microbiological Corrosion of Metallic Materials

549


40 Electronic Materials, Components, and Devices

559

R. P. Frankenthal and L. F. Garfias-Mesias

41 Corrosion of Electronics: Lead-Free Initiatives

565

M. Reid and L. F. Garfias-Mesias

42 Metastable Alloys

571

K. Hashimoto

43 Carbon Steel—Atmospheric Corrosion

579

I. Matsushima

44 Carbon Steel—Corrosion in Freshwaters

589

I. Matsushima

45 Carbon Steel—Corrosion by Seawater

601

I. Matsushima

46 Carbon Steel—Corrosion by Soils

609

I. Matsushima

47 Localized Corrosion of Iron and Steel

615

I. Matsushima

48 Weathering Steel

621

T. Murata

49 Corrosion of Steel in Concrete

633

J. P. Broomfield

50 Ethanol Stress Corrosion Cracking of Carbon Steels

649

J. Beavers, F. Gui, and N. Sridhar

51 Austenitic and Ferritic Stainless Steels

657

M. A. Streicher, Revised by J. F. Grubb

52 Duplex Stainless Steels

695

M.-L. Falkland, M. Glaes, and M. Liljas

53 Martensitic Stainless Steels

707

J. F. Grubb

54 Aluminum and Aluminum Alloys

715

E. Ghali

55 Cobalt Alloys P. Crook and W. L. Silence

747

CONTENTS

56

Copper and Copper Alloys

757

C. A. C. Sequeira

57

Lead and Lead Alloys

787

F. E. Goodwin

58

Magnesium and Magnesium Alloys

809

E. Ghali

59

Nickel and Nickel Alloys

837

D. C. Agarwal and N. Sridhar

60

Tin and Tinplate

853

T. P. Murphy

61

Titanium and Titanium Alloys

861

J. Been and J. S. Grauman

62

Zinc

879

X. G. Zhang

63

Zirconium Alloy Corrosion

893

B. Cox

PART IV 64

CORROSION PROTECTION

Controlling Flow Effects on Corrosion

903

K. D. Efird

65

Erosion–Corrosion: Recognition and Control

907

J. Postlethwaite and S. Nesˇic´

66

Using Plastics, Elastomers, and Composites for Corrosion Control

915

P. R. Khaladkar

67

Corrosion Control of Steel by Organic Coatings

971

C. H. Hare

68

Selection and Use of Coatings for Underground or Submersion Service

985

R. Norsworthy

69

Engineering of Cathodic Protection Systems

1001

J. H. Fitzgerald III

70

Stray-Current Analysis

1013

J. H. Fitzgerald III

71

Corrosion Inhibitors S. Papavinasam

1021

xiii

xiv

CONTENTS

72 Computer Technology for Corrosion Assessment and Control

1033

S. Srinivasan

PART V

TESTING FOR CORROSION RESISTANCE

73 Principles of Accelerated Corrosion Testing

1047

D. L. Jordan

74 High-Temperature Oxidation—Testing and Evaluation

1053

C. A. C. Sequeira

75 Testing for Flow Effects on Corrosion

1059

K. D. Efird

76 Accelerated Testing of Electronics to Simulate Long-Term Worldwide Environments

1063

L. F. Garfias-Mesias and M. Reid

77 Testing for Environmentally Assisted Cracking

1071

R. D. Kane

78 Test Methods for Wet H2S Cracking

1087

M. Elboujdaini

79 Atmospheric Corrosion Testing

1095

D. L. Jordan

80 Galvanic Corrosion Testing

1101

X. G. Zhang

81 Testing of Aluminum, Magnesium, and Their Alloys

1103

E. Ghali

82 Testing of Polymeric Materials for Corrosion Control

1107

B. Thomson and R. P. Campion

83 Corrosion Testing of Refractories and Ceramics

1117

M. Rigaud

84 Evaluation and Selection of Corrosion Inhibitors

1121

S. Papavinasam

85 Practical Corrosion Prediction Using Electrochemical Techniques

1129

D. C. Silverman

86 Electrochemical Noise D. A. Eden, Revised and updated by Q. J. Meng, M. Mendez, and M. Yunovich

1167

CONTENTS

PART VI 87

CORROSION MONITORING

Corrosion Monitoring

1181

P. R. Roberge

88

Diagnosing, Measuring, and Monitoring Microbiologically Influenced Corrosion

1203

B. J. Little, R. I. Ray, and J. S. Lee

Glossary of Selected Terms Used in Corrosion Science and Engineering

1217

Index

1223

xv

FOREWORD

In the roughly 10 years since the appearance of the second edition of the Corrosion Handbook, new technologies and new engineering systems have found their way into the global marketplace at an increasing rate. It is no wonder now that a third edition would be timely and appropriate. It is also no surprise that the third edition would expand the scope of the previous editions to include chapters on composites such as are used in airframes, shape memory alloys which find application in medical devices, and electrodeposited nanocrystals as well as application-specific chapters which address the materials of construction of the engineered barriers for nuclear waste containment, ethanol-induced stress corrosion cracking of carbon steels, and other such topics. An entirely new section on corrosion monitoring also appears in the third edition. Corrosion is ubiquitous: All engineering systems are subject to environmental degradation in service environments, whether these systems are used to meet the energy needs of the inhabitants of this planet; to provide clean air; to treat and transport water, food, and other products typical of our commercial world; to both save and improve the quality of our lives; and to ensure the readiness of those engineering systems that are of importance in terms of national defense and homeland security as well as many others. From heart stents to nuclear electric generating stations, corrosion is part of our world. The Corrosion Handbook continues today, as it has since its first appearance over 60 years ago, to serve as a trusted resource to generations of corrosion engineers. There are many reasons to believe that its presence in the libraries of engineering practitioners of all kinds is greater now than ever before. First, it appears that much of the expertise in this area of technology, which resided for decades in the staff and laboratories of metal producers, has retired and is not being

replaced as many of the metal producers have responded to the global economy of the past decade and more. Second, the interest of young people in engineering education, including corrosion engineering, is also in decline. Third, as the global economy recovers from the meltdown of the recent past, nations with a strong manufacturing base that creates products of value to the market will respond most quickly. But this will require an educated and informed engineering workforce. It is a concern to me that industrialized nations all over the world are on the brink of losing this technological infrastructure through retirement, the decline of traditional manufacturing industries, and declining student interest. Without a means of capturing this expertise in a useful form the next generation of engineers are going to find a gap in their knowledge base. I am confident that this volume will be of value in that context. Every industrialized nation must have the capacity and intellectual strength necessary to design, manufacture, and maintain either contemporary engineering systems or emerging engineering systems that may find their way into the marketplace of the future. The Corrosion Handbook remains an invaluable resource in that regard, and once again Winston Revie has assembled a world-class group of authors in producing a comprehensive volume covering the entire field of contemporary corrosion engineering. R. M. LATANISION February 2010 Director (Emeritus) The H. H. Uhlig Corrosion Laboratory, MIT, and Corporate Vice President Exponent–Failure Analysis Associates, Inc. Natick, Massachusetts

xvii

FOREWORD TO THE SECOND EDITION

The first and, prior to the current volume, only edition of the Corrosion Handbook was published in 1948. It represented a heroic effort by Professor Herbert Uhlig and the leadership of the relatively newly established Corrosion Division of the Electrochemical Society. It was intended, as Professor Uhlig recorded in the Preface to the 1948 edition, to serve as “. . . a convenient reference volume covering the entire field of corrosion, to bring together, in effect, much of the information scattered broadly throughout the scientific and engineering literature.” Its success was equally heroic: the Corrosion Handbook has served generation after generation of corrosion engineer and today, more than a one-half of a century since its first appearance, the volume remains a trusted resource in the personal libraries of many of those who populate the world’s engineering community. Over the years that I knew Professor Uhlig personally, he often mentioned to me his concern for the need to produce a revised edition of the Handbook. I am confident that he would have been very pleased that one of his doctoral students at MIT, Winston Revie, had taken up this challenge. Winston, just as his mentor, is a meticulous and innovative corrosion scientist. This truly monumental revision of the Corrosion Handbook is certain to serve the engineering community well as we enter the new millennium. Much has happened in corrosion science and engineering since 1948, and the contributors to this volume, an assembly of the international leaders in the field, have captured these changes wonderfully well. The breadth of corrosion and corrosion control is made

clear by the inclusion of ceramics, polymers, glass, concrete and other materials as well as of metals, the focus of the first edition. The introduction of standards into corrosion science and engineering is emphasized as is life prediction, and economic and risk analyses associated with environmental degradation of materials. While the introduction of new technologies has dramatically changed virtually every aspect of life on the Earth in the fifty years since the appearance of the Corrosion Handbook, what remains a persistent reality in the engineering enterprise is that engineering systems are built of materials. Whether an airframe, integrated circuit, bridge, prosthetic device or, perhaps as we shall see in the not too distant future, implantable drug delivery systems—the chemical stability of the materials of construction of such systems continues to be a key element in determining their useful life. This new edition of the Corrosion Handbook will serve, among others, designers, inspectors, owners and operators of engineering systems of all kinds, many of which are unknown today, for generations to come. Dr. Revie has succeeded, just as did his mentor in 1948, in producing a convenient reference volume covering the entire field of corrosion. R. M. LATANISION H. H. Uhlig Corrosion Laboratory Massachusetts Institute of Technology Cambridge, Massachusetts

xix

PREFACE

The objective in preparing this third edition of Uhlig’s Corrosion Handbook has been to provide an updated book—one affordable volume—in which the current state of knowledge on corrosion is summarized. The fundamental scientific aspects and engineering applications of new and traditional materials and corrosion control methods are discussed, along with indications of future trends. The book is intended to meet the needs of scientists, engineers, technologists, students, and all those who require an up-to-date source of corrosion knowledge. This new edition contains a total of 88 chapters divided among six parts: I. II. III. IV. V. VI.

Basics of Corrosion Science and Engineering Nonmetals Metals Corrosion Protection Testing for Corrosion Resistance Corrosion Monitoring

Topics discussed in chapters that are new in this edition include failure analysis (Chapter 1), principles of accelerated corrosion testing (Chapter 73), metal–matrix composites (Chapter 35), nanocrystals (Chapter 37), ethanol stress corrosion cracking (Chapter 50), computation of Pourbaix diagrams at elevated temperature (Chapter 9), high-temperature oxidation (Chapters 20 and 74), dealloying (Chapter 11), and diagnosing, measuring, and monitoring microbiologically

influenced corrosion (MIC) (Chapter 88). Dr. Tomomi Murata has provided some very insightful introductory notes on the effects of climate change, life-cycle design, and corrosion of steel under changing atmospheric conditions. Throughout the book, extensive reference lists are included to help readers identify sources of information beyond what could be included in this one-volume handbook. It is a pleasure to acknowledge the authors who wrote the chapters of this edition as well as the reviewers, who, in anonymity, carried out their work in the spirit of continuous improvement. I would also like to acknowledge the members of the Editorial Advisory Committee, who made many constructive suggestions to help define, focus, and clarify the discussions in this new edition. I would like to acknowledge Mary Yess and her staff at The Electrochemical Society Headquarters in Pennington, New Jersey, for their support during the preparation of this book. I greatly appreciate the encouragement and support of Bob Esposito and his staff at John Wiley & Sons, Inc. in Hoboken, New Jersey. Finally, I would like to thank my many friends and colleagues at the CANMET Materials Technology Laboratory, where it has been my privilege to work for the past 32 years. R. WINSTON REVIE Ottawa, Ontario, Canada

xxi

CONTRIBUTORS

Agarwal, D. C., DNV Columbus, Inc., Dublin, Ohio, USA

Bale, C. W., Departement de genie physique et de genie des materiaux, Ecole Polytechnique, Montreal, Quebec, Canada Beavers, J, A., DNV Columbus, Inc., Dublin, Ohio, USA Been, J., Alberta Innovates Technology Futures, Calgary, Alberta, Canada B€ ohni, H., Institute of Materials Chemistry and Corrosion, Swiss Federal Institute of Technology, Z€urich, Switzerland (Retired)

Fitzgerald III, J. H., Grosse Pointe Park, Michigan, USA Ford, T. E., University of New England, Biddeford, Maine USA Frankenthal, R. P., Bell Laboratories, Lucent Technologies, Murray Hill, New Jersey, USA (Retired) Garfias-Mesias, L. F., DNV Columbus, Inc., Dublin, Ohio, USA Ghali, E., Department of Mining, Metallurgy and Materials, Laval University, Quebec, Canada

Broomfield, J. P., Corrosion Consultant, London, UK

Glaes, M., Outokumpu Stainless AB, Avesta, Sweden

Campion, R. P., MERL Ltd., Wilbury Way, Hitchin, UK

Goodwin, F. E., International Lead Zinc Research Organization, Inc., Research Triangle Park, North Carolina, USA

Cox, B., Centre for Nuclear Engineering, University of Toronto, Toronto, Ontario, Canada (Retired) Crook, P., Haynes International, Kokomo, Indiana, USA (Retired)

Eden, D. A., Honeywell Process Solutions, Houston, Texas, USA

Eiselstein, L. E., Exponent-Failure Analysis Associates, Inc., Menlo Park, California, USA Efird, K. D., Efird Corrosion International, Inc., The Woodlands, Texas, USA Elboujdaini, M., CANMET Materials Technology Laboratory, Ottawa, Ontario, Canada Erb, U., Department of Materials Science and Engineering, University of Toronto, Toronto, Ontario, Canada Falkland, M. L., Outokumpu Stainless AB, Avesta, Sweden

Grambow, B., La Chantrerie, Laboratoire SUBATECH (UMR 6457), Ecole des Mines de Nantes, Nantes Cedex 3, France Grauman, J. S., TIMET, Henderson, Nevada, USA Grubb, J. F., Technical & Commercial Center, ATI Allegheny Ludlum Corp., Brackenridge, Pennsylvania, USA Gu, J.-D., School of Biological Science, The University of Hong Kong, Hong Kong, China Gui, F., DNV Columbus, Inc., Dublin, Ohio, USA Hare, C. H., Coating System Design Inc., Lakeville, Massachusetts, USA (Retired) Hashimoto, K., Tohoku Institute of Technology, Sendai, Japan

Deceased.

xxiii

xxiv

CONTRIBUTORS

Heidersbach, R., Dr. Rust, Inc., Cape Canaveral, Florida, USA

Mitton, D. B., Gold Standard Corrosion Science Group, LLC, Boston, Massachusetts, USA

Hihara, L. H., Department of Mechanical Engineering, University of Hawaii at Manoa, Honolulu, Hawaii, USA

Morris, P. I., FPInnovations, Vancouver, BC, Canada

Huet, R., Exponent-Failure Analysis Associates, Inc., Menlo Park, California, USA

Murata, T., Office of Technology Transfer Innovation Headquarters, Japan Science and Technology Agency, Tokyo, Japan

Jack, T. R., University of Calgary, Calgary, Alberta, Canada

John, R. C., Shell International E&P, Inc., Houston, Texas, USA Jordan, D. L., Ford Motor Company, Dearborn, Michigan, USA Kane, R. D., iCorrosion LLC, Houston, Texas, USA Kaye, M. H., Faculty of Energy Systems and Nuclear Science, University of Ontario Institute of Technology, Oshawa, Ontario, Canada Khaladkar, P. R., E. I. DuPont de Nemours & Co, Inc., Wilmington, Delaware, USA King, G. A., CSIRO Building, Construction and Engineering, Highett, Victoria, Australia (Retired) Kruger, J., Department of Materials Science and Engineering, Johns Hopkins University, Baltimore, Maryland, USA (Retired) Latanision, R. M., Exponent-Failure Analysis Associates, Inc., Natick, Massachusetts, USA

Murphy, T. P., Campion Hall, University of Oxford, Oxford, UK

Nesic, S., Institute for Corrosion and Multiphase Flow Technology, Ohio University, Athens, Ohio, USA Nessim, M., C-FER Technologies Inc., Edmonton, Alberta, Canada Norsworthy, R., Lone Star Corrosion Services, Lancaster, Texas, USA Papavinasam, S., CANMET Materials Technology Laboratory, Hamilton, Ontario, Canada

Parkins, R. N., University of Newcastle upon Tyne, Newcastle upon Tyne, UK

Pelton, A. D., Departement de genie physique et de genie des materiaux, Ecole Polytechnique, Montreal, Quebec, Canada Piro, M. H., Department of Chemistry and Chemical Engineering, Royal Military College of Canada, Kingston, Ontario, Canada

Lee, J. S., Naval Research Laboratory, Stennis Space Center, Mississippi, USA

Postlethwaite, J., Department of Chemical Engineering, University of Saskatchewan, Saskatoon, Saskatchewan, Canada (Retired)

Lewis, B. J., Department of Chemistry and Chemical Engineering, Royal Military College of Canada, Kingston, Ontario, Canada

Ray, R. I., Naval Research Laboratory, Stennis Space Center, Mississippi, USA

Liljas, M., Outokumpu Stainless AB, Avesta, Sweden

Rebak, R. B., GE Global Research, Niskayuna, New York, USA

Little, B. J., Naval Research Laboratory, Stennis Space Center, Mississippi, USA Malhotra, V. M., Consultant, Ottawa, Ontario, Canada

Reid, M., Stokes Research Institute, University of Limerick, Limerick, Ireland

Technology,

Revie, R. W., CANMET Materials Technology Laboratory, Ottawa, Ontario, Canada

Mendez, M., Honeywell Corrosion Solutions, Houston, Texas, USA

Rigaud, M., Departement de genie physique et de genie des materiaux, Ecole Polytechnique, Montreal, Quebec, Canada

Matsushima, I., Maebashi Maebashi, Japan

Institute

of

Meng, Q. J., Honeywell Corrosion Solutions, Houston, Texas, USA Mitchell, R., Laboratory of Microbial Ecology, Harvard School of Engineering and Applied Sciences, Harvard University, Cambridge, Massachusetts, USA

Deceased.

Roberge, P. R., Department of Chemistry and Chemical Engineering, Royal Military College of Canada, Kingston, Ontario, Canada Ruddick, J. N. R., Department of Wood Science, Forest Sciences Centre, University of British Columbia, Vancouver, B.C., Canada

CONTRIBUTORS

Sequeira, C. A. C., Instituto Superior Tecnico, Lisboa, Portugal Shibata, T., Department of Materials Science and Processing, Graduate School of Engineering, Osaka University, Japan (Retired)

xxv

Thompson, W. T., Centre for Research in Computational Thermochemistry, Royal Military College of Canada, Kingston, Ontario, Canada Thomson, B., MERL Ltd., Wilbury Way, Hitchin, UK

Silence, W. L., Consultant, Fairfield Glade, Tennessee, USA

Verink, Jr., E. D., Department of Materials Science and Engineering, University of Florida, Gainesville, Florida, USA (Retired)

Silverman, D. C., Argentum Solutions, Inc., Chesterfield, Missouri, USA

Wang, Y.-Z., Canadian Nuclear Safety Commission, Ottawa, Ontario, Canada

Sridhar, N., DNV Columbus, Inc., Dublin, Ohio, USA

Wilmott, M., Wasco Coatings Ltd., Kuala Lumpur, Malaysia

Srinivasan, S., Advanced Solutions–Americas, Honeywell International, Inc., Houston, Texas, USA

Young, A. L., Humberside Solutions Ltd., Toronto, Ontario, Canada

Staehle, R. W., University of Minnesota, Minneapolis and Industrial Consultant, North Oaks, Minnesota, USA

Yunovich, M., Honeywell Corrosion Solutions, Houston, Texas, USA

Streicher, M. A., E. I. DuPont de Nemours & Co., and the University of Delaware, Newark, Delaware

Deceased.

Zhang, X. G., Teck Metals Ltd., Mississauga, Ontario, Canada

INTRODUCTORY NOTES ON CLIMATE CHANGE, LIFE-CYCLE DESIGN, AND CORROSION OF STEEL T. MURATA Japan Science & Technology Agency, Saitama, Japan

A. CLIMATE CHANGE Climate change is attributed mainly to increased CO2 in our atmosphere because of anthropogenic activities and is expected to increase as much as 50% by 2030 compared to the concentration in 2005, that is, 359 ppm [1, 2]. Such a change will affect the corrosion of carbon steel through acidification due to increased concentration of HCO3 and Ca2þ in waters at temperatures a few degrees Celsius higher than those in 1990. In addition, other influential factors that will arise from climate change include the following: 1. Increase in precipitation 2. Formation of aerosols with CO2 emission 3. Increased SOx emissions caused by the use of sulfurbearing coal due to oil shortages 4. Enhanced biological growth in waters For these reasons, the corrosivity of environments in the future will be complex, and a simple acidification model will not be adequate. To predict the effects of climate change on corrosion, computational analyses and systematic corrosion studies are required to develop models based on projected climate change. “Time of wetness” is universally considered to be a key corrosion index for atmospheric corrosion. In recent years, weather instability has led to changes in global rainfall distribution, changes that could lead to new and different predictive indices for atmospheric corrosion. For corrosion in waters, microbiological factors are expected to increase in importance with the changing climate. In contrast to the environmental factors that pertain to corrosion in air and water, the heterogeneous distribution of chemicals in

contaminated soils in industrialized areas results in nonuniform soil corrosivity. Dynamic corrosion models are required with on-site monitoring systems.

B. LIFE-CYCLE DESIGN To minimize the environmental burden and to attain a sustainable society, life-cycle design of steel structures is required to ensure safety, reliability, durability, and the best use of materials and energy throughout the life cycle. The life-cycle concept will be required for future design and construction of social as well as industrial infrastructure. For example, in developing a life-cycle design for weathering steels, discussed in Chapter 48, reliable corrosion data for long-term service and a systematic approach to minimize both corrosion damage and social costs are necessary. In general, corrosion is studied using a set of parameters under simplified or fixed conditions. In the real world, in response to constantly changing environmental parameters, corrosion behavior also changes. For this reason, an understanding of corrosion dynamics is required, and the corrosion protection models that are implemented must have a capacity to reflect dynamic environmental conditions that are subject to constant change.

REFERENCES 1. Intergovernmental Panel on Climate Change (IPCC), “Climate and Water,” Technical Paper VI, Geneva, June 2008. 2. Intergovernmental Panel on Climate Change (IPCC), “Implications of Proposed CO2 Emissions Limitations,” Technical Paper IV, Geneva, Oct. 1997, Figure 6, p. 16.

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THE ELECTROCHEMICAL SOCIETY SERIES

Corrosion Handbook Edited by Herbert H. Uhlig Modern Electroplating, Third Edition Edited by Frederick A. Lowenheim Modern Electroplating, Fifth Edition Edited by Mordechay Schlesinger and Milan Paunovic The Electron Microprobe Edited by T. D. McKinley, K. F. J. Heinrich, and D. B. Wittry Chemical Physics of Ionic Solutions Edited by B. E. Conway and R. G. Barradas High-Temperature Materials and Technology Edited by Ivor E. Campbell and Edwin M. Sherwood Alkaline Storage Batteries S. Uno Falk and Alvin J. Salkind The Primary Battery (in Two Volumes) Volume I Edited by George W. Heise and N. Corey Cahoon Volume II Edited by N. Corey Cahoon and George W. Heise Zinc-Silver Oxide Batteries Edited by Arthur Fleischer and J. J. Lander Lead-Acid Batteries Hans Bode Translated by R. J. Brodd and Karl V. Kordesch Thin Films-Interdiffusion and Reactions Edited by J. M. Poate, M. N. Tu, and J. W. Mayer Lithium Battery Technology Edited by H. V. Venkatasetty Quality and Reliability Methods for Primary Batteries P. Bro and S. C. Levy

Techniques for Characterization of Electrodes and Electrochemical Processes Edited by Ravi Varma and J. R. Selman Electrochemical Oxygen Technology Kim Kinoshita Synthetic Diamond: Emerging CVD Science and Technology Edited by Karl E. Spear and John P. Dismukes Corrosion of Stainless Steels A. John Sedriks Semiconductor Wafer Bonding: Science and Technology Q.-Y. Tong and U. G€oscle Uhlig’s Corrosion Handbook, Second Edition Edited by R. Winston Revie Atmospheric Corrosion Christofer Leygraf and Thomas Graedel Electrochemical Systems, Third Edition John Newman and Karen E. Thomas-Alyea Fundamentals of Electrochemistry, Second Edition V. S. Bagotsky Fundamentals of Electrochemical Deposition, Second Edition Milan Paunovic and Mordechay Schlesinger Electrochemical Impedance Spectroscopy Mark E. Orazem and Bernard Tribollet Fuel Cells: Problems and Solutions Vladimir S. Bagotsky

PART I BASICS OF CORROSION SCIENCE AND ENGINEERING

1 CORROSION FAILURE ANALYSIS WITH CASE HISTORIES L. E. EISELSTEIN AND R. HUET Exponent-Failure Analysis Associates, Inc., Menlo Park, California

A. Introduction B. Failure analysis procedures C. General approach for conducting a failure analysis C1. Gather general information C2. Formulate hypotheses C3. Gather further information C4. Validate or reject hypotheses D. Techniques typically used to investigate corrosion failures D1. Sampling and collecting corrosion evidence D2. Determining corrosion rates D3. Characterizing the form of corrosion D3.1. Uniform corrosion D3.1.1. Corrosion rates D3.1.2. Cosmetic corrosion failure D3.2. Galvanic corrosion D3.3. Crevice corrosion D3.4. Pitting D3.5. Dealloying D3.6. Erosion–corrosion D3.7. Environmentally assisted cracking D4. Complex failure analysis References

A. INTRODUCTION Are failure analyses useful? The answer is an emphatic Yes. There are many reasons to perform a failure analysis; the most common one is to help prevent future failures. To ensure that corrective actions will be effective, it is necessary to understand why failures have occurred in the first place. Otherwise, any design or manufacturing changes that are implemented may not be effective or simpler ways to prevent

future failures may be overlooked. Another common reason for performing a failure analysis is to establish responsibilities for the mishap. For instance, an insurance company may want to determine if an event is covered by the policy or not or financial responsibilities must be established for the resolution of a lawsuit. There is always value in performing some level of failure analysis, even if it seems that it would be better to try something new rather than finding out exactly what went wrong. A tremendous amount of information can be gleaned from understanding how things fail, and this knowledge is invaluable in making things (equipment, machines, and processes) work better in the future. There is a natural tendency to move past the setback of a failure, maybe to avoid dwelling on unpleasant facts or assuming that nothing can be learned from something that did not work out. However, much can be learned from understanding what went wrong, and a good learning opportunity should not be thrown away with the failed parts. The end point of a failure analysis depends on the specific circumstances and the type of answers needed. In some cases, it may be enough to rule out a specific failure mode, rather than establishing exactly what happened. In an industrial setting, the goal may be to understand the failure enough to be able to identify corrective measures or to determine if a product recall is required. Finally, in serious accidents an exhaustive failure analysis may be necessary for insurance, legal, or safety reasons.

B. FAILURE ANALYSIS PROCEDURES The analysis of corrosion failures is not fundamentally different from any other failure analysis. Although some

Uhlig’s Corrosion Handbook, Third-Edition, Edited by R. Winston Revie Copyright 2011 John Wiley & Sons, Inc.

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CORROSION FAILURE ANALYSIS WITH CASE HISTORIES

unique techniques may be used and the failure modes are specifically related to corrosion, the methodology used is much the same for every type of failure analysis. Failure analysis follows the scientific method. Typically there is a question to be answered, such as “How or why did this failure occur?” Hypotheses are proposed to answer the question or questions. The hypotheses are checked against facts, experiments, and analyses. In the end, some hypotheses are ruled out and others are confirmed. In practice, there may be many hypotheses put forward to explain a failure, some likely and others very unlikely. Conducting an efficient failure analysis means that most effort is spent proving the hypothesis that ultimately turns out to be the correct explanation while dismissing early (but with good reason) those hypotheses that turn out to be unfounded. Thus it is important to identify the relevant information early on, even though one cannot be sure of what information is ultimately going to be relevant. Performing a failure analysis results in an interplay between hunches, developing likely hypotheses, testing them rigorously, dismissing other hypotheses for good cause, and keeping an open mind for other possible scenarios if the ones that appeared likely at first turn out to not fit all of the facts.

C. GENERAL APPROACH FOR CONDUCTING A FAILURE ANALYSIS There is no firm set of rules to conduct a correct failure analysis, but the following approach will help. Gather some general information, formulate hypotheses, and then use these hypotheses to gather more targeted information. Use this targeted information from observations, testing, and analyses to validate or rule out the hypotheses. C1. Gather General Information The first step of a failure analysis should be to understand the role of the failed component and its environment. Is the failed part available for examination? Does a cursory examination provide some clues as to why it failed? By definition, the failure was not desired, so were any steps taken in the design or operation to prevent it? Was the part or equipment that failed a recent design, had it been modified recently, or did it have a long history of good service? Does the manufacturer, designer, operator, eyewitnesses, or end user have any hypotheses about the cause of this failure? All these questions will help orient the investigator at the start of a failure analysis. C2. Formulate Hypotheses It is important to formulate hypotheses early in a failure investigation, because they will guide the collection of further information. Without some hypotheses, relevant information

may be overlooked, or to the contrary too much information will be gathered in an effort to be inclusive, which may also impair getting to the truly important facts. One should also think about hypotheses that have to be considered even if they may be ruled out in the end. For instance, if the corrosion failure involves dissimilar metals, galvanic corrosion should be investigated because it is an obvious possibility, even if other aspects of the situation make it unlikely. C3. Gather Further Information The hypotheses will help the investigator gather relevant information. It is important to collect facts that may tend to disprove a given hypothesis as well as those that may support it. One of the traps to be avoided is to bias the information collected toward proving one particular scenario. The Royal Society’s motto Nullius in verba, roughly translated as “Take nobody’s word for it,” is a valuable principle to follow when gathering information. Reported observations, hearsay (“Joe told me that. . .”), and sweeping generalizations (“This has never happened before”) should be noted but not considered reliable until they have been checked for accuracy. Often the simple act of verifying information will separate fact from fiction and considerably clarify a picture that may have appeared confused at first. C4. Validate or Reject Hypotheses It is very important that hypotheses be tested or validated in some fashion; without this step they are nothing more than speculation. The validation may be very simple in some cases or it may require extensive analysis in other cases, but this step should never be overlooked. The validation process must be based on physical and engineering principles, not merely on a process of elimination based on commonalities and differences. Sometimes a failure analysis proceeds by listing common factors and differences between failures and instances of successful operation. This method may be useful as a guide to formulate hypotheses (although not as useful as trying to understand the physical factors affecting a failure), but it should not be used as the exclusive means to validate or reject any hypothesis. The validation of hypotheses must not be biased in favor of a specific scenario. It is not always easy to recognize that a hypothesis that looked promising at first should actually be rejected or modified, but one should remain alert to this possibility if the validation does not turn out as expected.

D. TECHNIQUES TYPICALLY USED TO INVESTIGATE CORROSION FAILURES Corrosion failures often involve the use of some specific information or techniques. Several guidance documents

TECHNIQUES TYPICALLY USED TO INVESTIGATE CORROSION FAILURES

suggest checklists or procedures that are specific to corrosion failures, for instance the American Society for Testing and Materials (ASTM) has issued a Standard Guide for Corrosion-Related Failure Analysis [1]. These guides supplement the general failure analysis process and adapt it to corrosion failures. A few topics that recur regularly in corrosion failures are discussed here, but the reader may wish to consult some of the extensive literature on failure analysis [1–14]. D1. Sampling and Collecting Corrosion Evidence Sampling and collection of corrosion evidence are often key steps in corrosion failure analysis. Sampling should be done carefully because corrosion products and deposits often contain valuable information that can be easily damaged or contaminated. For example, if microbial activity is suspected, samples should be collected in sterile containers under appropriate conditions to avoid contamination by other microorganisms. These live samples should be analyzed promptly. If the pH or dissolved oxygen level of the aqueous environment may help to explain the corrosion, field measurements should be made of these parameters as they can change rapidly during storage. Samples should be protected from contamination by other debris and stored such that they will be protected from further corrosion damage. In addition, samples must be representative. In many cases a few well-chosen samples will be enough. The investigator may easily choose a few “typical” samples from the affected and nonaffected areas. However, in cases where one must assess the condition of a large number of items (e.g., to assess the extent of damage), a statistically valid sampling method must be used. There is no sampling method that always produces a “statistically valid” sample, so the sampling must be defined for each case with the help of a competent statistician. D2. Determining Corrosion Rates A question often arises about the timing of some corrosion event. For instance, one may want to know for how long some corrosive conditions have been present or how long before some corrosion will result in a leak or vessel rupture. Unfortunately, corrosion rates are notoriously variable, and in some cases it is nearly impossible to make precise predictions. However, corrosion rates have been published for many combinations of materials and environment and they can be useful if their limitations are understood. These rates are typically averages over many samples or observations; although the performance of any single sample may deviate significantly from the average, in aggregate these rates can be useful. In general, corrosion rates tend to slow down with time, so it is usually important to know the time period over which the rate was measured because linear extrapolation cannot be used in most cases.

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Direct measurements of corrosion rates in the laboratory usually take a long time, which may not be practical in the context of the failure analysis. Where this is not possible, accelerated corrosion tests are sometimes performed; however, it is generally quite difficult to determine the acceleration factor over the actual environment. When possible, long-term monitoring of the corrosion process in service is the best way to obtain relevant rates, and this monitoring should be started as soon as possible after the equipment, product, or process is placed in service. D3. Characterizing the Form of Corrosion Various types of corrosion have been defined, based generally on morphology or mechanism. Authors differ on this classification, but generally some variation of the following eight categories of corrosion are used: (1) uniform, (2) galvanic, (3) crevice, (4) pitting or localized, (5) intergranular, (6) dealloying, (7) erosion–corrosion, (8) environmentally assisted cracking (EAC), or stress corrosion cracking (SCC) [15]. Other named forms of corrosion such as microbiological-induced corrosion (MIC), filiform corrosion, and liquid-metal embrittlement are typically covered as subcategories of these types. In this chapter we discuss corrosion failures of various types and illustrate them with case histories where applicable. D3.1. Uniform Corrosion. Uniform corrosion, also known as general corrosion, is a very common corrosion type where the metal is corroding more or less uniformly. Most often, this type of corrosion is easily investigated, since it occurs whenever a susceptible metal is in contact with an aggressive environment: Leave a nail in water and it will corrode. Questions that are more difficult to answer may include the rate at which the corrosion damage occurred, for instance, why the corrosion was particularly rapid in a specific case. In other cases, the key may be to find out how the environment came into contact with the susceptible metal. Finally, questions may arise regarding the effect of corrosion on the strength or other properties of the corroding material. D3.1.1. Corrosion Rates. Although corrosion rates are extremely variable, they are useful to indicate the average behavior of many samples. They are also useful as a broad indicator of the intensity of corrosion: In cases where the observed rates are far different from the published ones, there should be some explanation of the difference. Case Study: Corrosion Rates Indicate That Product Is Not Suitable for Intended Use. A company decided to manufacture a copper-covered stainless sheet to make roof panels, gutters, and flashing products. In this process the stainless steel was electroplated with copper on both sides: 10 mm on the side expected to be exposed to the weather and

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FIGURE 1.1. Copper-coated stainless steel (thick side).

3 mm on the other (see Fig. 1.1). Atmospheric exposure tests performed by the manufacturer and others indicated a corrosion rate of the order of 1 mm per year for coastal marine atmospheric environments. At this rate, the copper coating would have been consumed in about 10 years (or 3 years if the wrong side of the sheet is exposed to atmospheric corrosion), well short of the 30-year life that was contemplated for this product and somewhat shorter than the time to form a patina. Once the copper layer is removed, the roof looks like stainless steel rather than the intended copper patina, and furthermore the now-exposed stainless steel may pit. Although this material may perform well in dry and noncoastal marine climates, it was clearly not suitable for general use under all outdoor atmospheric corrosion conditions. The field experience indicated that after less than one year of service there were complaints of excessive corrosion. The copper layer was completely removed from areas of severe exposure, such as in chimney flashings exposed to the acidic flue gases or from the water runoff from wood shakes (see Fig. 1.2). Even though most of the installations had not failed after a few years, the occurrence of several early failures corroborated the reported corrosion rates and indicated that the coatings were too thin for this application. By comparison, a similar product with 50 mm of copper roll bonded to both sides of stainless steel sheet has demonstrated good performance for more than 20 years of service. Case Study: Extraordinarily High Corrosion Rates. Dilute nitric acid can be extremely corrosive to carbon steel, but concentrated nitric acid passivates carbon steel. This passivation is temporary and can be reversed. A wellknown experiment illustrates this behavior: A nail is placed in a test tube and concentrated nitric acid is added to cover about half the nail. Nothing happens because the nail is

FIGURE 1.2. Rapid corrosion of copper coating from areas underneath wood shake. Arrows indicate bare stainless steel exposed after the copper layer has corroded away.

passivated by the acid. Water is added slowly so that it forms a separate layer on top of the nitric acid. Nothing happens at first, but after a minute or so, corrosion starts at the interface where water dilutes the nitric acid. The passivation breaks down and the corrosion reaction becomes extremely violent. This experiment was repeated unwittingly inside a nitrogen tetroxide (N2O4) tank car through a series of errors [16]. The carbon steel tank car was used to carry N2O4 that was being used in a paper plant. The N2O4 will react with water to form nitric acid. This N2O4 tank car had been involved in an earlier incident in which a significant amount of water had entered it undetected. Sometime later, the presence of water was detected and it was decided to drain the car of the nitrogen tetroxide and nitric acid that had formed. After draining some material, the car was erroneously thought to be nearly empty and water was added to dilute what was thought to be a small “heel” of liquid but was in fact a significant amount of nitric acid and nitrogen tetroxide. The operation was repeated twice in the following days, every time with the same result. Each time, a complicated process of mixing and diluting was taking place inside the car: The inflow of water may have stirred the car’s content, but water, concentrated nitric acid, and nitrogen tetroxide tended to separate in layers because of their different densities. The result was unpredictable and at some point the carbon steel passivation broke down. Very rapid corrosion occurred, a massive release of nitrogen oxides overwhelmed the venting capacity of the pressure relief valve, and the tank car ruptured. There were three distinct corrosion bands on the inside of the car where severe loss of material had occurred in a matter of hours or days. D3.1.2. Cosmetic Corrosion Failure. Sometimes, uniform corrosion does not affect the structural properties of the corroding part and the only effects are cosmetic. Whether or not this represents a failure depends on the circumstances.


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If the appearance of the part is of no concern, a slight corrosion may be acceptable; otherwise corrosion may be deemed a serious concern well before any structural failure may occur. Case Study: Atmospheric Corrosion on the Underside of Roofing Panels. The roof of a California building was made of steel panels with a thick coating on the external side. The underside, which was exposed to an attic space, was left with a thin shop primer coating because there were no concerns about its appearance. After a few years in service, the structural performance of the panels was called into question, in part because the underside had visibly corroded. Some atmospheric corrosion had appeared in spite of the shop primer. Cross sections of the corroded areas demonstrated that the depth of corrosion was minimal and that, at the observed rate, perforation of the roofing panel would not occur during the expected lifetime of the building. Consequently, this was an instance where corrosion was not a failure. Case Study: Atmospheric Corrosion of Terne-Coated Roof. In two buildings, one located in Alaska and the other in Louisiana, roofs made from lead–tin (terne)– coated stainless steel became severely discolored instead of developing the dull gray appearance typically associated with lead roofs. Metals such as copper and lead used for roofs develop a patina or surface film on exposure to the atmosphere; the patina for terne is usually dull gray. However, in these two instances, the roofs developed irregular patterns of reddish-yellowish corrosion patterns (Fig. 1.3). Our investigation showed that the discoloration of the terne-coated roofs was not caused by rusting of the stainless steel substrate. Rather, it resulted from the normal patination of the terne coating being disrupted due to adverse environmental conditions. The sequence of patina formation is orthorhombic lead oxide (PbO, yellow) ! basic lead carbonate ! normal lead carbonate ! lead sulfite ! lead sulfate. If access to the air is restricted, there may not be enough carbon dioxide to form the basic lead carbonate, so the yellow lead oxides remain. There was no attack of the stainless steel substrate and no risk of leaks from the roof corrosion, so this was purely a cosmetic failure. However, this is an example where the appearance of the roof was important—the terne coating had been chosen for its color— so this was indeed a corrosion failure. D3.2. Galvanic Corrosion. Galvanic corrosion is a common failure mode. It occurs where a less noble metal is in electrical contact with a more noble metal in an electrolyte. ASTM standard G71 provides a test procedure to evaluate the potential for galvanic corrosion [17]. A related failure mode, but not strictly galvanic corrosion, occurs when the source of potential difference results from

FIGURE 1.3. Terne-coated roof showing severe reddish-yellowish discoloration instead of the expected gray patina.

a difference in the electrolyte composition between two zones. For instance, differences in oxygen concentration can accelerate corrosion of the area depleted in oxygen. Yet another source of potential differences may be electrical currents generated by some external cause, for instance, stray ground currents from large electrical equipment. Typically, galvanic corrosion results in fairly rapid attack. But potential differences may be used to protect a piece of metal by forcing the corrosion to occur on a sacrificial anode. This, of course, is the principle of cathodic protection. However, the cathodic protection may not perform as anticipated, leading to corrosion failures. Case History: Ineffective Cathodic Protection. In some small ships, the propeller shafts are enclosed in stern tubes that are part of the hull. The inside of the stern tubes is exposed to the seawater whereas the outside is dry, being in the hold of the ship. In a specific model of ships, the propeller shafts are made of stainless steel while the stern tubes and the hull are made of low-carbon steel. The shafts are supported by three bearings in the tubes: one aft, one midlength of the tubes, and one that is part of the forward end of the tubes. The hull is painted and there are zinc anodes on the hull and propeller for cathodic protection. However, severe corrosion was observed after less than one year in service on the inside of the stern tubes, particularly at their forward end, which is the farthest from the aft opening of the tubes (Fig. 1.4). The corrosion pattern was typical of painted low-carbon steel exposed to aerated seawater without cathodic protection. Potential measurements on a boat in service confirmed that the steel was not cathodically protected in the forward region of the stern tube. This is not surprising: There were no zinc anodes inside the tubes and the “throwing power” of the anodes on the outside of the hull was simply not enough to reach the inner surface at the forward end of the tubes, especially given the shaft bearing supports that are in the way. Further, the stainless steel shaft, sealing flange, and

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a three-parameter Weibull distribution (Fig. 1.5) whereas the homes without recirculation could not be easily characterized without accounting for home location or hot water usage.

FIGURE 1.4. Severe corrosion of the inside of a stern tube. The pattern of general corrosion with some deeper pits is typical of lowcarbon steel in seawater.

rotating seal were likely to have anodically polarized the steel tubes, accelerating the corrosion. Under these conditions, any small defect or holiday in the paint coat (and the inside surfaces of the tubes are difficult to paint) would lead to the observed corrosion. It was recommended that zinc anodes should be placed on the inside of the tubes for effective cathodic protection. Case History: Thermogalvanic Corrosion. A relatively uncommon source of potential difference between two electrodes is a difference in temperature. The electrode potential of a metal piece in an electrolyte depends in part on the temperature, but temperature effects are usually small and negligible. However, in some cases these temperature differences lead to significant failures. A common occurrence is the potential difference between the hot and cold water lines in homes when the copper pipes are buried under a concrete slab. In several residential developments in the western United States, the hot and cold water lines were buried under the home cement slab foundation. The temperature difference between the hot and cold copper water lines provides a potential difference, which in some environments results in the hot line being anodic to the cold. The soil and soil moisture provide the electrolyte. The result is pitting corrosion on the outside of the hot water lines. This failure mode happens more often on homes with hot recirculation pumps. These pumps keep hot water circulating constantly in the lines, so that there is almost instantly hot water at the faucets when they are open. The result is that the lines are hot 24 hours a day, whereas without recirculation pumps the lines would be cold most of the time. In this situation, statistical methods can be used to predict future leaks based on prior experience. The Weibull distribution is commonly used to model failures. In a particular development, the number of leaks in homes with recirculation pumps could be modeled accurately by

D3.3. Crevice Corrosion. Crevice corrosion is a type of localized corrosion at an area that is shielded from full exposure to the environment. This type of attack is usually associated with small volumes of stagnant solution caused by holes, gasket surfaces, lap joints, surface deposits, and crevices under bolt and rivet heads. This form of corrosion is sometimes also called deposit or gasket corrosion [15]. ASTM provides a guide for evaluation of the crevice corrosion resistance of stainless steel and nickel-based corrosionresistant alloys in chloride-containing environments [18]. Crevices are formed in a variety of design situations. For instance, aircraft skin panels are joined by lap splice joints, where moisture and corrosive atmospheric gases can become trapped between the two panels, resulting in crevice corrosion [19, 20]. This mechanism was blamed for the 1988 Aloha Airline accident in which a 20-year-old Boeing 737 lost a major portion of the upper fuselage at 25,000 ft [19, 21, 22]. In this particular instance, crevice corrosion resulted in the formation of voluminous hydrated aluminum oxides which acted to separate the two skin panels, stressing the rivets and resulting in fatigue failures. D3.4. Pitting. Alloys that maintain their resistance to corrosion through the formation of a protective passive layer, such as stainless steels, aluminum alloys, and titanium alloys, generally do not suffer from uniform corrosion; rather they will usually corrode as a result of the localized breakdown of a small region of the passive film. Corrosion occurs rapidly at this defect compared to the surrounding material covered with the passive film, resulting in the formation of a pit. ASTM has a standard that helps to characterize the nature of pitting [23]. Case History: Pitting Corrosion of Aluminum Due to Copper-Containing Fungicides. Irrigation pipes used in several California farms suffered from rapid pitting corrosion starting on the inside of the pipe. The pipes are used to irrigate crops and spray fungicide and other chemicals. The pipes that experienced severe pitting were used to spray copper-based fungicides. A literature review indicates that dissolved copper (in the form of copper hydroxide, the active ingredient of the fungicides) is not enough to promote severe pitting of aluminum alloys; there must be some chlorine and bicarbonate ions as well. The pitting mechanism has not been identified in detail, but it probably involves penetration of the aluminum oxide layer by chloride ions, plating of copper on the exposed aluminum surface, and starting the formation of corrosion products nodules over the corroding areas. Once

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o c c u r r e n c e

90 80 70 60 50 40 30 20 10 5 2

2 2

1 –5

C D F %

o c c u r r e n c e

W/mle/t0ˆ t0 (1) = 278.7

2

C D F

–2 –1 –05 –1 Eta Beta n/s 8079 1. 447 616/766

–02

YF MC

%

–01 10

100

1000

10000

100000

Datum (units)

90 80 70 60 50 40 30 20

W/mle/t0ˆ t0 (1) = 717.9

10 5 2 1 –5 –2 –1 –05 –1 Eta 1591

–02

Beta 1. 723

n/s 4822/4564

YF MC

–01 10

100

1000

10000

100000

Datum (units)

FIGURE 1.5. Three-parameter Weibull distribution fits very well the cumlative failure distribution of homes with recirculation pumps (left), but homes without pumps cannot be characterized with a twoor three-parameter Weibull distribution without accounting for geographic distribution or other effects such as hot water usage (right).

the nodules are formed, they set up concentration cells and severe pits grow under the nodules. Experiments confirmed that all three ions were necessary for severe pitting. Aluminum pipes exposed to fungicide dissolved in deionized water were slightly attacked, with few shallow pits forming. Similarly, pipes exposed to farm water (containing bicarbonate hardness as well as about 9 ppm chlorides) without fungicide experienced mild pitting. Only the combination of farm water and fungicides caused severe pitting, comparable to the damage observed in the irrigation pipes on the farms (Fig. 1.6). D3.5. Dealloying. Dealloying (also known as selective dissolution) includes two commonly occurring phenomena known as graphitic corrosion of cast irons and dezincification of brasses. These are two related corrosion modes, in which

one of the components of an alloy leaches out selectively, leaving behind a spongy and weak matrix. In dezincification, zinc leaches out of the brass, leaving behind a copper matrix. In graphitic corrosion of cast irons, the iron corrodes away, leaving behind a porous and weak graphite matrix. This dealloying corrosion may not be easily recognized by a simple visual examination; frequently the matrix does not look appreciably different from the intact material. A metallographic cross section may be required to identify this type of corrosion and determine its extent. Case History: Graphitic Corrosion of a Sewer. A fire protection water main at a restaurant failed. As part of the repairs, slurry comprised of cement and pea gravel was pumped under the restaurant foundation to fill the void that was created from water gushing out of the ruptured

FIGURE 1.6. Pitting observed on the inside surface of an aluminum irrigation pipe (left) and cross-section of a pit in a laboratory sample of aluminum pipe exposed to copper-containing fungicide (right).

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fire main. The contractor performing the repairs then found that the sewer line was plugged with this slurry. Further inspection found the sewer line was corroded and the contractor thought that perhaps his cement had corroded the pipe. The insurance for the restaurant ordered an investigation into the cause of the sewer line failure. The piping is shown in Figure 1.7. It was immediately clear that pumping the slurry was not likely to have corroded the sewer line, as cement is not corrosive to ferritic materials. Inspection of the gray cast iron sewer line showed severe graphitic corrosion: substantial wall thinning, holes, and selective leaching of iron from the cast iron matrix. The corrosive attack was primarily on the sewer pipe inner surface, with little or no corrosive attack from the outside. The graphitic corrosion was extensive, indicating that the sewer may have been leaking for a long time, probably years. The most severe corrosion, including holes in the sewer line, was near the foundation on the south side of the restaurant where the fire protection water main entered the building. Examination of the fire main showed that it had failed due to weakening of the pipe wall from external corrosion. Thus, it is likely that the sewer line leak was the source of moisture that caused the external corrosion damage to the fire protection water main. D3.6. Erosion–Corrosion. Some metals corrode more rapidly when exposed to flowing water. A good example is copper: Corrosion is much accelerated when exposed to water flowing at more than about 1 m/s past the surface. These conditions may be found in home plumbing systems with hot recirculation lines, as shown below. Case History: Erosion–Corrosion of Copper Hot Recirculation Lines. Some homes have hot recirculation lines; that is, the hot water lines form a loop in which a small pump keeps a constant flow of water. This allows hot water to flow almost instantly from any tap in the home. However, the constant flow and high temperature could lead to severe corrosion. In a recent residential development where these loops were installed in fairly small homes, the available pumps were oversized with respect to the

pipes, resulting in a constant high-velocity flow of hot water. Localized corrosion developed in a few years at areas where the flow was disrupted, such as elbows and tees. The corrosion pattern showed deep, well-defined eroded areas looking like sand dunes or river erosion patterns, indicative of erosion–corrosion (Fig. 1.8). D3.7. Environmentally Assisted Cracking. EAC/SCC is a frequent cause of corrosion failures because it tends to be unexpected. Many early corrosion failures occurred in situations where SCC was unknown or unexpected at the time. For instance, the Silver Bridge at Point Pleasant, West Virginia, was designed in 1927 at a time when SCC was not known to occur under rural atmospheric conditions in the classes of bridge steels used for construction [24–26]. Yet SCC developed in an eyebar suspension link and led to the catastrophic collapse of the bridge after 40 years in service. Similarly, the Flixborough explosion of 1974, which killed 28 people in Great Britain when 50 tons of cyclohexane was released in a chemical plant, was (in part) a result of nitrate SCC and liquid metal embrittlement [27, 28]. Generally, SCC requires a specific combination of stress, material, and environment. Since the fundamental mechanisms for SCC are not always well understood, new combinations that can cause SCC are unfortunately found by accident. Case History: Room Temperature Transgranular SCC of Austenitic Stainless Steel. Austenitic stainless steel, such as type 304 or 316, is known to be susceptible to transgranular SCC (TGSCC), but it was thought that this required exposure to concentrated chloride environments above 60 C. However, there have been several recent reports of such SCC at room temperature. For instance, SCC has been observed in permanent anchors used for rock climbing in Thailand (Fig. 1.9), in limestone formations at sea level [29]. After a few years, some of the anchors developed extensive cracking, to the point that they broke under normal climbing loads. Examination of the broken pieces showed

FIGURE 1.7. Fire main and sewer piping from restaurant (left) and graphitic corrosion on inner surface of sewer piping (right).


11

FIGURE 1.8. Copper hot recirculation line showing signs of erosion–corrosion.

pervasive intergranular SCC and the fracture surfaces contained a high concentration of magnesium, rather than the sodium that would be expected in a marine environment. Review of the known cases of room temperature transgranular SCC showed that the climbing anchors are in an environment that combines known promoters of TGSCC. Essentially, this type of SCC requires that the steel be exposed to very high concentrations of chlorides, which are promoted by salts such as magnesium or calcium chlorides, and by low relative humidity (a high relative humidity tends to dilute any salt water film on the metal). The climbing cliffs are located on tower karsts, which get their characteristic steep sides because of very active dissolution and redeposition of the limestone. Thus the climbing anchors can be exposed to calcium and magnesium salts as well as sodium chloride from the ocean. Although the environment is usually very humid, there are times where the relative humidity is fairly low, concentrating any solution that has formed on the stainless steel. Thus the climbing anchors are exposed to an environment that is extremely severe for transgranular chloride SCC.

D4. Complex failure analysis Sometimes the incident under investigation is particularly complex or involves significant costs, either human or financial. In these instances, it is more important than ever to perform the failure analysis in a careful manner, separating various issues that may be involved and thoroughly validating hypotheses before coming to final conclusions. Such an example is described here. Case Study: Chlorine Release at a Manufacturing Facility. A massive amount of chlorine gas was released to the atmosphere at a chemical plant making chlorine from calcium chloride [30, 31]. In the plant, the chlorine gas is liquefied before transport by railcar. In the liquefaction process, the chlorine gas is first compressed, then cooled down in a shell-and-tube heat exchanger, in which the chlorine flows inside tubes while chilled calcium chloride brine at about 23 C ( 10 F) circulates on the outside of the tubes. The liquefied chlorine is sent to a storage tank via a long transfer pipe containing several tees and elbows. The release

FIGURE 1.9. Typical 316L rock climbing bolt (left) installed permanently on climbing routes in cliff formations located on Thai beaches (right).

12


occurred through a large corrosion hole in an elbow in the transfer line between the liquefier and the storage tank. Early indications showed that the hole in the elbow was due to severe general corrosion on the inside surface. Dry chlorine is not corrosive to the steel elbow, but addition of water to the chlorine could make the mixture extremely corrosive. Potential sources of water included the chlorine liquefier: If there was a leak in the tubes, brine solution could enter the chlorine stream. The liquefier was pressure tested and several tubes were found to be leaking. The liquefier was cut open to expose the leaking tubes. To the general surprise, an old rag was found stuck in the shell of the liquefier, right at the brine inlet piping. It appeared that the rag had been there for a long time, most probably from the time of installation of the liquefier, some 25 years earlier. The rag partially blocked the brine flow path, resulting in accelerated flow in the areas that remained unobstructed. The leaking tubes were found in the area of accelerated brine flow. With these early findings, the following scenario was hypothesized: The rag had been in the liquefier since installation, leading to increased brine flow rates over some tubes. This eroded the tubes, and, after about 25 years, one or more tubes were perforated. The brine flowed into the chlorine, creating a very corrosive mixture that corroded through the elbow in the transfer line within a few days, before the water contamination of the chlorine could be detected. Each step in the scenario was validated with tests and analyses. 1. The rag increased the flow rate over some tubes in the bundle. A fluid flow finite-element analysis of the liquefier inlet section was performed. With the cloth, the brine velocity in the area of the holes was about 3.94 m/s, whereas without the cloth it would have been only 1.27 m/s, or about one-third. 2. This increased flow rate led to through erosion of the tube in about 25 years. The observed corrosion rate of the tubes in areas where the brine flow was not accelerated by the cloth was very slow, about 10 mm per year. At this rate, it would take over 200 years to puncture a tube, so the design and operation of the liquefier were not the cause of the tube leak. The corrosion rate of the tube material exposed to brine flowing at high velocities was measured in a test bed in which chilled brine of various pH values was flowed over dummy tubes at various velocities. For a brine solution at its natural pH (no chemical additions), flowing at about 4 m/s, the interpolated corrosion rate would lead to through-wall erosion in about 22 years, very close to the actual service life of 25 years. 3. Once the tube was perforated, brine flowed into the chlorine stream. This conclusion was actually not immediate, because both the brine and chlorine

systems were pressurized to roughly the same value. A careful pressure drop analysis of both the chlorine and the brine systems, coupled with review of plant data and some actual measurements made on a mockup of the brine system, confirmed that the brine pressure was likely higher than that of the chlorine, so that brine would be entrained into the chlorine stream. 4. The transfer line elbow corroded mostly from the inside out. This was not immediately obvious, as both the inside and the outside of the elbow were severely corroded. The outside surface may have suffered from underinsulation corrosion during its life, and perhaps more rapid corrosion had occurred during the incident, when liquid chlorine was released from the leak site, but the extent of damage on the inside and outside needed to be compared accurately. A cross section of the failed elbow was traced over the outline of an undamaged elbow; this illustrated clearly that most of the corrosion had occurred on the inside surface (Fig. 1.10). 5. The corrosion rate at the elbow was high enough to lead to a leak in a few days. It is likely that the brine leak in the chlorine liquefier occurred relatively shortly before the incident, because it is likely that small amounts of water in the chlorine would have been detected prior to distribution. The best estimate of the start of the brine leak was a few days before the incident. The elbow was about 5.3 mm thick originally, so the corrosion rate would have to be extraordinarily high for through-wall attack in a few days. A related observation that required explanation was that the corrosion at the leak site was much more severe than

FIGURE 1.10. Erosion–corrosion wall loss of liquefied chlorine run down elbow compared to exemplar elbow.

REFERENCES

13

FIGURE 1.11. CFD calculated shear rate of (a) failed elbow-tee configuration and (b) elbow closest to storage tank without tee.

elsewhere in the transfer line, even though the line contained at least one other elbow. This was investigated by a combination of tests and analysis. The failed elbow happened to be located immediately after a T fitting, so that the flow went through two changes of direction in rapid succession. A computer fluid dynamics (CFD) model showed that the flow velocity and the shear rate in this elbow were much higher than in the other elbow in the transfer line, explaining why the corrosion was so much more severe at the leak point (Fig. 1.11). Actual corrosion rate measurements with several mixtures of brine and chlorine, both static and flowing, showed very high corrosion rates, but not quite high enough to achieve penetration in less than seven days. However, given the variability of the measured rates and the uncertainty about the flow conditions and the composition of the corroding mixture, it was concluded that the elbow most likely corroded in a few days after the tubes had started leaking. The hypothesized scenario was thus validated step by step. This was a very serious incident, in which the stakes were high enough that a thorough failure analysis had to be performed, and the resources were available. Most failure analyses cannot be this detailed, but the basic steps must be the same: gather initial information, formulate hypotheses, and, most importantly, validate these hypotheses with the level of detail appropriate for the conclusions to be reached.

3.

4.

5. 6. 7.

8.

9. 10.

11.

12.

REFERENCES 1. American Society for Testing and Materials (ASTM), ASTM G161-00(2006): Standard Guide for Corrosion-Related Failure Analysis, ASTM, West Conshohocken, PA, 2006. 2. American Society for Testing and Materials (ASTM), ASTM E2332-04: Standard Practice for Investigation and Analysis of

13.

14. 15.

Physical Component Failures, ASTM, West Conshohocken, PA, 2004. American Society for Testing and Materials (ASTM), ASTM E1492-05: Standard Practice for Receiving, Documenting, Storing, and Retrieving Evidence in a Forensic Science Laboratory, ASTM, West Conshohocken, PA, 2005. American Society for Testing and Materials (ASTM), ASTM E1459-92(2005): Standard Guide for Physical Evidence Labeling and Related Documentation, ASTM, West Conshohocken, PA, 2005. Journal of Engineering Failure Analysis, Elsevier Ltd., Oxford, UK, 2009. Journal of Failure Analysis and Prevention, ASM International, Materials Park, OH, 2009. P. de Castro and A. Fernandes, “Methodologies for Failure Analysis: A Critical Survey,” Mater. Design, 25(2), 117–123 (2004). H. Herro and R. Port, The Nalco Guide to Cooling Water System Failure Analysis, McGraw-Hill Professional, New York, 1992. E. D.D. During, Corrosion Atlas: A Collection of Illustrated Case Histories, Third Edition, Elsevier, Amsterdam, 1997. American Society for Testing and Materials (ASTM), ASTM E860-07: Standard Practice for Examining and Preparing Items That Are or May Become Involved in Criminal or Civil Litigation, ASTM, West Conshohocken, PA, 2007. American Society for Testing and Materials (ASTM), ASTM E1188-05: Standard Practice for Collection and Preservation of Information and Physical Items by a Technical Investigator, ASTM, West Conshohocken, PA, 2005. K. Esaklul, Handbook of Case Histories in Failure Analysis, Vols. I and II, ASM International, Materials Park, OH, 1993. ASM, Metals Handbook, Vol. 11, Failure Analysis and Prevention, 11th ed., ASM International, Materials Park, OH, 2002. C. Brooks and A. Choudhury, Failure Analysis of Engineering Materials, McGraw-Hill Professional, New York, 2001. M. G. Fontana and N. D. Greene, Corrosion Engineering, McGraw-Hill Book Company, New York, 1978.

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16. National Transportation Safety Board (NTSB), Tank Car Failure and Release of Poisonous and Corrosive Vapors, Gaylord Chemical Corporation, Bogalusa, Louisiana, Oct. 23, 1995NTSB Report No. HZB-98-01, Washington, D.C. 1998. 17. American Society for Testing and Materials (ASTM), ASTM G71-81 (2009): Standard Guide for Conducting and Evaluating Galvanic Corrosion Tests in Electrolytes, ASTM, West Conshohocken, PA, 2007. 18. American Society for Testing and Materials (ASTM), ASTM G78-01(2007): Standard Guide for Crevice Corrosion Testing of Iron-Base and Nickel-Base Stainless Alloys in Seawater and Other Chloride-Containing Aqueous Environments, ASTM, West Conshohocken, PA, 2007. 19. R. G., Kelly, “Crevice Corrosion,” in ASM Handbook: Corrosion Fundamentals, Testing, and Protection, ASM International, Metals Park, OH, 2003, pp. 242–247. 20. W. Wallace, D. Hoeppner,and P. Kandachar, AGARD Corrosion Handbook, Vol. 1, Aircraft Corrosion: Causes and Case Histories, Advisory Group for Aerospace Research and Development, Neuilly sur Seine, France, 1985. 21. National Transportation Safety Board (NTSB), Aircraft Accident Report Aloha Airlines, flight 243, Boeing 737-200, N73711, near Maui, Hawaii April 28, 1988, NTSB No. AAR89/03, Washington, D.C. 1989. 22. P. R. Roberge, Handbook of Corrosion Engineering, McGrawHill, New York, 2000. 23. American Society for Testing and Materials (ASTM), ASTM G46-94(2005) Standard Guide for Examination and

24.

25. 26.

27.

28.

29.

30.

31.

Evaluation of Pitting Corrosion, ASTM, West Conshohocken, PA, 2005. J. Bennett and H. Mindlin, “Metallurgical Aspects of the Failure of the Point Pleasant Bridge,” J. Test. Eval. JTEVA, 1(2), 152–161 (1973). C. LeRose, “The Collapse of the Silver Bridge,” West Virginia Historical Soc. Quart., Charleston, WV, 15(4) (2001). National Transportation Safety Board, (NTSB), Collapse of U.S. 35 Highway Bridge Point Pleasant, West Virginia December 15, 1967, NTSB No. HAR-71/01, Washington, D.C. 1970. S. Mannan and, F., Lees, “Flixborough,” in Lees’ Loss Prevention in the Process Industries: Hazard Identification, Assessment, and Control, Vol. 3, Elsevier, Amsterdam, 2005. A. Cottrell and P. Swann, “Technical Lessons of Flixborough— A Metallurgical Examination of the 8-Inch Line,” Chem. Eng., 308, 266–274 (1976). A. Sjong and L. E. Eiselstein, “Marine Atmospheric SCC of Unsensitized Stainless Steel Rock Climbing Protection,” J. Failure Anal. Prevention, 8(5) (2008). S. P. Andrew et al. “Evaluation of a Failure in a Chlorine Production Facility,” in Proceedings of IMEC2001: 2001 ASME International Mechanical Engineering Congress and Exposition, ASME, New York, 2001. J. G. Routley, “Massive Leak of Liquified Chlorine Gas Henderson, Nevada (May 6, 1991),” Report 052, FEMA— United States Fire Administration, Washington, DC, 1991.

2 COST OF METALLIC CORROSION J. KRUGER† Department of Materials Science and Engineering, Johns Hopkins University, Baltimore, Maryland

A. Introduction B. Economic effects of metallic corrosion in the United States B1. Approach B1.1. Elements of corrosion costs used in the BCL I/O model B1.2. Scenarios for the I/O model B1.3. The I/O model B2. Results B3. 1995 Update of the NBS–BCL study C. Corrosion costs in various countries C1. United Kingdom C2. German Federal Republic C3. Sweden C4. Finland C5. Union of Soviet Socialist Republics (now Russia) C6. Australia C7. India C8. Japan C9. Previous U.S. Studies D. Cost of selected corrosion control measures References

A. INTRODUCTION Metallic corrosion seriously affects many sectors of a nation’s economy or, on a vastly smaller scale, the design choices made by an engineer. This finding is so because corrosion and protective measures to control corrosion result in the utilization of materials, energy, labor, and technical expertise that would otherwise be available for alternative uses. Corrosion causes users of metal products to incur added expenses such as †

more costly corrosion-resistant materials, painting and other corrosion-protective measures, earlier replacement of capital goods, increased spare-parts inventories, and increased maintenance. Some of these costs are avoidable and could be lowered by applying the economically best available corrosion prevention technology. However, decreasing the presently remaining unavoidable costs would require advances in technology. The cost of the corrosion of metals will be considered from two standpoints: the cost to the economy of a nation and the cost of selected corrosion control measures. The main basis for the discussion of the cost to a nation will be the study published in 1978 [1] that the U.S. Congress directed the National Bureau of Standards (NBS; currently National Institute of Standards and Technology, NIST) to undertake. This study was, and still remains, probably the most comprehensive investigation of the full extent of corrosion on the economy of a nation. The analysis required was contracted out to Battelle Columbus Laboratories (BCL). The NBS–BCL study, unlike previous ones (B3), was based upon a solid technical–economic method that attempted to evaluate, in a rigorous way, all costs of corrosion, direct and indirect, over the entire economy and to evaluate the uncertainties in these cost estimates. “A significant feature of the study was that the method employed—input/output analysis—provides a methodological framework that permits comprehensive treatment of all elements of the costs of corrosion: production costs, capital costs, and changes in useful lives, for example. The input/output model allows analysis of interindustry relationships in the national economy and attribution of relative costs to specific segments of the economy” [1a, p. 30]. To carry out the second aspect of this discussion, a National Association of Corrosion Engineers (NACE) Recommended Practice [2] and a more recent NACE update [3] will be the source of the discussion of the economic appraisal of a selected corrosion control method.

Retired.

Uhlig’s Corrosion Handbook, Third Edition, Edited by R. Winston Revie Copyright Ó 2011 John Wiley & Sons, Inc.

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COST OF METALLIC CORROSION

B. ECONOMIC EFFECTS OF METALLIC CORROSION IN THE UNITED STATES The NBS–BCL study [1] of metallic corrosion costs in the United States in 1975 employed a modified version of the BCL National Input/Output (I/O) model. A series of articles giving a detailed discussion of the NBS–BCL study [4] was published in the more available corrosion literature. B1. Approach B1.1. Elements of Corrosion Costs Used in the BCL I/O Model. The BCL I/O model using 130 sectors of the U.S. economy included the following as costs associated with corrosion: (a) material and labor expenditures associated with the protective measures of painting, applying cathodic protection, coatings, and use of inhibitors; (b) the expenses arising from the extra material and labor for prevention; (c) the partial corrosion losses that result in replacement costs and lost production; and (d) the expenses incurred by using information, technology transfer, research, development, and demonstration of methods to cope with the destructive effects of metallic corrosion. A basic factor to be considered in measuring macrocorrosion costs in a sector of an economy was the lifetime and replacement value of a given component that could suffer from corrosion damage. Table 2.1 gives some of the elements of the costs of corrosion used in the BCL I/O model. Corrosion costs used in the I/O approach can be either direct or indirect. The NBS report [1a, p. 11] describes these in the following way: Direct costs include all reductions in the requirements for inputs for production which would become possible if there were no corrosion. These include flow inputs (e.g., pig iron into steel), capital inputs for expansion and replacement of capacity (e.g., blast furnaces for steel), and TABLE 2.1. Some Elements of the Costs of Corrosiona Capital costs Replacement of equipment and buildings Excess capacity Redundant equipment Control costs Maintenance and repair Corrosion control Design costs Materials of construction Corrosion allowance Special processing Associated costs Loss of product Technical support Insurance Parts and equipment inventory a

See [1a, p. 10].

value added. The direct flow effects include reduced maintenance costs and the use of less expensive materials for embodiment in outputs. Among the direct capital effects are the reduced need for equipment due to less time down for maintenance and the lower replacement cost because of increased equipment life. The value-added effects include reduced costs of labor and lower depreciation allowances for the smaller capital requirements. The indirect effects include two elements: (1) In addition to the reduced input requirements, the inputs also cost less, because of savings in their own and earlier production processes: (2) the general interactive effects of reductions in production levels on one another. Value added is the additional value accruing to a product’s ingredients as they are fashioned into the product itself. It includes wages, salaries, rents, profits, interest, taxes, and depreciation. It can also be defined as the value of the productive factors contributed by the industry itself, rather than those purchased from other industries. B1.2. Scenarios for the I/O Model. Three scenarios were developed and used in the I/O model to quantify corrosion costs: Three “worlds” were formulated: world I—the realworld economy in 1975, world II—an imaginary corrosionfree world, and world III—a hypothetical world in which everyone applies the best economically practical corrosion control measures. The scenarios used in the model involved the gross national product (GNP—a more recently used term is gross domestic product, GDP). Three sets of these worlds were (a) world II minus world I is the total national cost of corrosion and represents resources that are wasted because of corrosion; (b) world III minus world I is the total national avoidable costs of corrosion and represents resources that would be available if economically best preventive practices were used throughout the economy; and (c) world III minus world II measures presently unavoidable costs. B1.3. The I/O Model. To describe the BCL I/O model, it is best to quote directly from the NBS report [1a, p. 12]: “A number of characteristics make input–output analysis, pioneered by W. W. Leontief, and the modified Battelle model well suited for use in estimating the total direct and indirect costs of corrosion. The model is quite detailed. In this study, it has 130 economic sectors and each is represented by a production function consisting of the respective inputs from that sector plus value added. As a result, relatively detailed industry corrosion cost data may be incorporated into the model for simulation purposes. The complex structure serves as a guide for the precise analysis of corrosion costs and a means for integrating the results.” “The model is comprehensive. It has sufficient components to allow all the contributions to corrosion costs (production expenses, capital cost, reductions in replacement, and excess capital capacity, for example) to be considered in the analysis. Because of the model’s structure, all of these

ECONOMIC EFFECTS OF METALLIC CORROSION IN THE UNITED STATES

aspects, and their interactions, may be evaluated in a coordinated and systematic manner.” “The model is simultaneous and, therefore, able to account for both direct and indirect effects of certain changes in the economy. This is critical to estimating the total costs of corrosion to society.” “Because the model simultaneously determines equilibrium values, comparative static analysis (i.e., comparison of alternate growth scenarios at the same moment in time) is an obvious application. For example, the costs of corrosion in the existing world (World I) are compared to those in each of the two hypothetical worlds mentioned previously— World II in which no corrosion exists and, thus, the costs are zero, and World III in which ‘best practice’ corrosion control methods are employed.” B2. Results If we use the approach outlined above, the NBS–BCL study separates the total costs of corrosion into two costs: (a) avoidable costs, which are costs that can be reduced “by the most economically effective use of presently available corrosion technology,” and (b) unavoidable costs, which result from “presently unavoidable losses.” For the study’s base year 1975, total costs of metallic corrosion (materials, labor, energy, and technical capabilities) were estimated by the BCL I/O model used in the study to be $82 billion, 4.9% of the $1677 billion GNP. Approximately 40% of this ($33 billion, 2.0% of the GNP) was estimated to be avoidable. Combining the BCL results and the NBS detailed analysis of the uncertainty, the total national yearly cost of metallic corrosion was reported to be about $70 billion (4.2% of GNP), with an uncertainty of 30%. The NBS report found the 1975 avoidable cost of corrosion was roughly 15% of the total, but it estimated that it could have a range of 10–45%. An analysis of the errors, especially the estimates of the avoidable costs, led to the conclusion in the NBS report that the values cited above were reasonable. Table 8 of the NBS report [1a] lists the total costs that were allocated to the 130 economic sectors of the United States that were produced by the BCL I/O model. This table gives corrosion losses on a dollar basis and as cost per unit of sales for the total and avoidable costs of both direct and direct plus indirect costs. In addition to the 130 sectors of the economy provided by Table 8 of this report, the following special area costs were covered in more detail in the NBS report: (a) U.S. federal government—total costs attributable to corrosion were estimated to be 2% of the federal budget with 20% of this total being judged to be avoidable. (b) Personally owned automobiles—total expenses of corrosion in the ownership of an automobile were

17

found to be $6–14 billion, 1975 dollars, with avoidable costs being $2–8 billion. (c) Electric power—total direct costs arising from the effects of corrosion on the operation of power generation plants were estimated to be $4.1 billion with avoidable costs of $120 million (3%of totalcosts). (d) Fossil fuel energy and materials—total costs arising from the additional energy and materials losses resulting from corrosion were estimated to be $1.4 billion, with avoidable costs of $248.5 million for energy, and $1.705 billion, with avoidable costs of $212 million, for nonrenewable raw material sectors. B3. 1995 Update of the NBS–BCL Study The Specialty Steel Industry of North America engaged Battelle to produce a report [5] updating the NBS–BCL study [1, 3] and reflecting the changes resulting from economic growth, inflation, and 20 years of scientific research and technological advances. In Table 2.2, the revised estimates for 1995 are compared to the 1975 BCL values—not taking the NBS uncertainty analysis into consideration. In part, the Battelle panel that produced the report attributed the 1995 reductions in the percent of GNP to the following factors: (a) the anticorrosion technology of the motor vehicle industry (the most significant factor); (b) increased use of stainless steels, coated metals, and more protective coatings; (c) substitution of material to reduce weight; and (d) reclassification of unavoidable costs as avoidable. In a more recent study, the annual direct cost of corrosion to industry and to governments in the United States was estimated to be approximately $276 billion, or 3.1% of the GDP [5a]. TABLE 2.2.

Metallic Corrosion in the United Statesa 1975 (billions of current dollars)

1995

All Industries Total Avoidable

82.0 33.0

296.0 104.0

Motor Vehicles Total Avoidable

31.4 23.1

94.0 65.0

Aircraft Total Avoidable

3.0 6

13.0 3.0

Other Industries Total Avoidable a

See [5].

47.6 9.3

189.0 36.0

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C. CORROSION COSTS IN VARIOUS COUNTRIES The fact that many countries have attempted to assess their national corrosion costs points up the worldwide awareness that corrosion can be a serious economic concern. It is useful to compare the United States results to those obtained in various industrial countries where such assessments of national corrosion costs were undertaken (including other past U.S. studies). These other studies of corrosion losses either involved major data gathering and interpretation efforts, which were the results of the analysis by a single authority of the country’s corrosion costs in terms of the knowledge of the best presently available corrosion control measures and industrial practices, or, simply, assertions by an author with no reference as to how the costs were computed. This is the rationale for using the NBS–BCL study as the major focus of this discussion, because it looks at the entire economy of a nation and seeks to estimate the uncertainties in the numbers reported. Because of this more rigorous approach, the results of the NBS–BCL analysis gave higher total costs of corrosion but, surprisingly, agreed qualitatively with those found in previous studies as discussed in the NBS report. In 1986, a more recent examination of corrosion costs in various countries was the topic of an NACE symposium entitled “International Approaches to Reducing Corrosion Costs” [6]. The NBS study considered the following countries.

TABLE 2.3. National Cost of Corrosion and Corrosion in the United Kingdoma Industry or Agency Building and construction General engineering Marine Metal refining and semifabrication Oil and chemical Power Transport Water Total a

Estimated Cost (£M) 250 11 280 15 180 60 350 25 1365

See [7].

found that some £310 million, or 23% of this total figure, was potentially avoidable. The estimated potential savings were £310 million, or 22.7% of the GNP. They suggested approaches toward achieving these savings, such as improved materials selection, specification and control of protective measures, improved awareness of corrosion, especially in design, and greater use of cathodic protection. An informal conference was held in 1971 as a supplement to the Hoar report to discuss the findings. Six sessions were held, one for each section of the report [8]. C2. German Federal Republic

C1. United Kingdom In 1969, a major data-gathering and interpretation endeavor, which was the precursor to the NBS–BCL U.S. study and unlike the less meticulous efforts of the countries described below, was initiated in the United Kingdom with the appointment by the Minister of Technology of a 25-member committee headed by one of the leading corrosionists in the United Kingdom, T. P. Hoar, to determine the cost of corrosion [7]. The committee contacted 800 industries in the country, all government departments, corrosion protection companies, and corrosion consultants. They were to gather from these sources information on the effects of corrosion, including the amount of shutdowns, rejection of product losses, structural failures, and the loss to industries from these. The committee added to these losses the costs of items replaced because of corrosion, expenditures on corrosion protection, and information services, research, and development in the various industries. Using these collected data, the Hoar committee arrived at an industrywide estimate of the cost of corrosion [7]. The Hoar report reported losses to the United Kingdom (Table 2.3) of £1.365 billion ($3.2 billion, 1969 U.S. dollars) for 1969–1970. This amounts to 3.5% of the GNP of the United Kingdom for that period. In addition, the committee

Behrens [9] estimated that total losses for the period 1968–1969 were 19 billion DM ($6 billion, 1969 dollars), with avoidable costs of 4.3 billion DM ($1.5 billion, 1969 dollars). No details were given as to what these figures include or how they were computed. Total costs were reported to be about 3% of the West German GNP for 1969, and avoidable losses were roughly 25% of total costs. These figures, with respect to GNP and percentage of avoidable cost, are in good agreement with figures found for other nations. C3. Sweden A partial study of corrosion costs in Sweden by Tr€adgaidh [10] in which painting expenditures to combat corrosion were analyzed for the year 1964 found these costs to be 300–400 million crowns ($58–77 million, 1964 dollars) with between 25 and 35% being avoidable. C4. Finland Costs to Finland for the year 1965 have been estimated by Vl€asaari [11] to be 150–200 million markaa ($47–62 million, 1965 dollars). Linderborg [12], referring to these losses, described the factors that must be taken into account in

REFERENCES

assessing corrosion costs to the Finnish nation. He recognized that an important factor was the variable lifetimes for a variety of items using the specific example of the automobile. C5. Union of Soviet Socialist Republics (Now Russia) Kolotyrkin [13] reported in 1969 that corrosion costs were 2% of the GNP, or 6 billion rubles ($6.7 billion, 1969 U.S. dollars), giving no indication as to what this figure includes or how it was computed.

19

$5.5 billion for the late 1940s. A more recent study was carried out by NACE [18]. This study, based on replies of 1006 persons to a questionnaire, estimated the cost to NACE members of direct expenditures of corrosion control measures to be $9.67 billion for 1975.

D. COST OF SELECTED CORROSION CONTROL MEASURES

For the period 1960–1961 Rajagopalan [15] estimated the cost of corrosion to India was 1.54 billion rupees ($320 million, 1961 dollars). He calculated the expenditures of certain measures to prevent or control corrosion, including direct material and labor expenses for protection, additional costs for increased corrosion resistance or redundancy, costs of information transfer, and funds spent on research and development. No quantitative estimate of uncertainty was attempted nor were avoidable and unavoidable costs broken down.

In order for the corrosion engineer to select the economically optimum corrosion control measure for a given specific problem, it is necessary that the corrosion engineer recognize the pertinent economic factors that bear on the choice of appropriate corrosion technology. To address this need, NACE Technical Unit Committee T-3C on Economics of Corrosion produced and issued in 1972 the Recommended Practice NACE Standard RP-02-72, “Direct Calculation of Economic Appraisals of Corrosion Control Measures” [2]. In 1994, the NACE Task Group T-3C-1 issued a technical report [3], “Economics of Corrosion,” to replace, the simpler and less rigorous from an accounting practice standpoint, RP-02-72. The objectives of the 1994 report were(1) present the economic techniques in a form that can be readily understood and used by engineers as a decision-making tool; (2) facilitate the communication of decisions between the corrosion technologist and management; and (3) justify investments in anticorrosion methods that have long-term benefits. The 1994 report used more advanced and standardized accounting notation and terminology that was based on American National Standards Institute (ANSI) standard Z94.5 entitled “Engineering Economy” [19]. Another more accessible source of the calculational techniques that enable an economic appraisal of corrosion control measures has been published by Verink [20]. A useful feature of the report (also present in the 1972 document) was a section devoted to worked examples and applications using the calculational techniques in the report to select process equipment and the best alternate cathodic protection proposal.

C8. Japan

REFERENCES

C6. Australia The direct costs of corrosion in 1973 were estimated by Revie and Uhlig [14] to be A$470 million ($550 million, 1973 dollars). The authors decided that these costs are “probably too low” considering the factors they used to develop this figure. Some additional direct costs—mostly labor—for the mining, transportation, and communications industries were unavailable. Only muffler corrosion was considered as contributing to automobile losses. Lifetimes were not taken into account quantitatively. The amount of $470 million was 1.5% of Australia’s GNP for 1973. However, since indirect costs may equal or exceed this figure, total corrosion costs to Australia were estimated to be 3% of GNP. No quantitative effort was made to assess uncertainties or to separate these costs into avoidable and unavoidable components. C7. India

A survey [16] conducted from 1976 to 1977 in Japan found that the annual direct cost of corrosion was 2500 billion yen ($9.2 billion, 1974 dollars), which amounts to 1.8% of the Japanese GNP. If the indirect costs were included, the total would increase severalfold. C9. Previous U.S. Studies Probably the first itemized measure of the costs of corrosion in the United States was carried out by Uhlig [17], who arrived at a value for the total direct corrosion losses of

1. L. H. Bennett, J. Kruger, R. I. Parker, E. Passaglia, C. Reimann, A. W. Ruff, H. Yakowitz, and E. B. Berman, “Economic Effects of Metallic Corrosion in the United States—AThree Part Study for Congress.” (a) Part I, NBS Special Publication 511-1, SD Stock No. 003-003-01926-7. (b) Part II, NBS Special Publication 511-2 Appendix B. A Report to NBS by Battelle Columbus Laboratories, SD Stock No. 003-003-01927-5, U.S. Government Printing Office, Washington, DC, 1978. (c) Part III Appendix C, Battelle Columbus Input/Output Tables, NBS GCR78-122, PB-279 430, National Technical Information Service, Springfield, VA, 1978.

20


2. NACE Standard RP-02-72, “Recommended Practice Direct Calculation of Economic Appraisals of Corrosion Control Measures,” National Association of Corrosion Engineers, Houston, TX, 1972. 3. NACE Technical Committee Report, Task Group T-3C-1 on Industrial Economic Calculational Techniques, “Economics of Corrosion,” National Association of Corrosion Engineers, Houston, TX, 1994. 4. J. H. Payer, W. K. Boyd, D. G. Dippold, and W. H. Fisher, Mater. Perform., 19(1–7) (1980). 5. Report by Battelle to Specialty Steel Industry of North America, “Economic Effects of Metallic Corrosion in the United States—A 1995 Update,” Apr. 1995. (a) G. H. Koch, M. P. H. Brongers, N. G. Thompson, Y. Paul Virmani, and J. H. Payer, Corrosion Costs and Preventive Strategies in the United States, Supplement to Materials Performance, July 2002, Report No. FHWA-RD-01-156, Federal Highway Administration, McLean, VA, 2002. See also www. corrosioncost.com. 6. NACE Symposium, “International Approaches to Reducing Corrosion Costs,” R. N. Parkins (Ed.), NACE, Houston, TX, March 1986. 7. T. P. Hoar, “Report of the Committee on Corrosion and Protection,” Department of Trade and Industry, H.M.S.O., London, U.K. 1971.

8. T. P. Hoar, Information Conference “Corrosion and Protection,” presented at the Instn. Mech. Engrs., April 20–21, 1971, to discuss the Hoar Committee. 9. D. Behrens, Br. Corros. J., 10(3), 122 (1975). 10. K. F. Tr€adgaidh, Tekn. Tedskrift (Sweden), 95(43), 1191 (1965). (quoted by Linderborg [12]). 11. V. Vl€asaari, Talousel€am€a’ (Economy), No. 14/15, 351(1965) (quoted by Payer et al. [4]). 12. S. Linderborg, Kemian Teollusius (Finland), 24(3), 234 (1967). 13. Y. Kolotyrkin, quoted in Sov. Life, 9, 168 (1970). 14. R. W. Revie and H. H. Uhlig, J. Inst. Engr. Aust., 46(3–4), 3 (1974). 15. K. S. Rajagopalan, Report on Metallic Corrosion and Its Prevention in India, CSIR. Summary published in “The Hindu,” English language newspaper (Madras), Nov. 12, 1973. 16. Committee on Corrosion and Protection, Boshoku Gijutsu (Corrosion Eng.), 26(7), 401 (1977). 17. H. H. Uhlig, Corrosion, 6(1), 29 (1950). 18. NACE Committee Survey Report, Corrosion, 31 (10) 1975. 19. ANSI Standard Z94.5 (latest revision), “Engineering Economy,” American National Standards Institute, New York. 20. E. D. Verink, “Corrosion Economic Calculations,” in ASM Handbook, Vol. 13, 9th ed., Corrosion, ASM International, Materials Park, OH, 1987, pp. 369–374.

3 ECONOMICS OF CORROSION* E. D. VERINK, JR.† Distinguished Service Professor Emeritus, Department of Materials Science and Engineering, University of Florida, Gainesville, Florida

A. Introduction B. Corrosion—an economic problem B1. Basics—money and time C. Notation and terminology D. Methods of economic analysis D1. Annual versus continuous compounding D1.1. First iteration E. Depreciation E1. Straight-line method E2. Declining-balance method E3. Declining balance switching to straight line E4. Sum-of-the-year’s-digits method F5. Accelerated cost recovery system method E6. Alternate cost recovery system F. Generalized equations G. Worked examples and applications References Bibliography

A. INTRODUCTION Over $220 billion is lost to corrosion in the United States each year, according to government and industry studies. This corrosion cost is equivalent to 3 or 4% of the gross

* Adapted from “Economics of Corrosion,” Publication 3C194, Technical Committee Report, copyright Ó 1994 by NACE International. Reprinted with permission. Further reproduction is prohibited. NACE technical committee reports are reviewed every 10 years; users should contact NACE at P.O. Box 218340, Houston, Texas 77218-8340. † Retired.

national product (GNP). The real tragedy of this annual corrosion cost is that 15% or more could be saved by the application of existing technology to prevent and control corrosion. The existing technology includes the following methodologies to prevent and control corrosion: proper design, selection of materials, coatings and linings, cathodic protection, and inhibitors. Standards, reports, books, and thousands of technical articles attest to the successful use of the existing technology to prevent and control corrosion and thus to reduce the annual losses to corrosion.

B. CORROSION—AN ECONOMIC PROBLEM Corrosion is essentially an economic problem. It is vitally important that engineers and engineering managers be aware of the economic impact their decisions have on the ability of a business to meet its corporate goals. A discipline that assists in the measurement of the economic impact of such decisions is called “engineering economy.” Engineering economy is concerned with money, both as a resource and as the price of other resources. Business success is dependent on the prudent and efficient use of all resources, including money. The principles of engineering economy permit direct comparisons of potential alternatives in monetary terms. In this way, they encourage efficient use of resources. The economics of corrosion evaluation involves the assessment of the technical validity and economic justification of each alternative. Corrosion technologists generally spend much effort exploring the technical validity of materials or processes. However, each corrosion problem may have more than one material or process that could satisfactorily solve the technical problem. Each candidate material or process probably has a unique stream of investment and operating


21

22

ECONOMICS OF CORROSION

or maintenance costs while providing equivalent technical benefits. Indeed, once the technically viable alternatives are chosen and their respective costs and performance characteristics are isolated, the technologist’s decision becomes one of financial analysis because almost all of the factors in the decision process can be reduced to the magnitude and timing of cash flows. The purposes of this chapter are: 1. To present the economic techniques in a form that can be readily understood and used by engineers as a decision-making tool 2. To facilitate the communication of decisions between the corrosion technologist and management 3. To justify investment in anticorrosion methods that have long-term benefits

B1. Basics—Money and Time Consider the effects of time and earning power on $20.00. If $20.00 is placed in the bank and earns interest at a rate of 5% per year, it grows to $21.00 when the interest, $1.00, is paid at the end of the year. Thus, $20.00 today at 5% is equivalent to $21.00 a year from now. Stated another way, in order to have $21.00 one year from now, only $20.00 has to be deposited today if the interest rate is 5%. The $20.00 is called the “discounted present value” of the $21.00 needed one year hence. The initial deposit as well as the earned interest left in the account have earning power because the interest is compounded, which means that it is computed on both the principal and the accrued interest. Suppose the $20.00 were used to pay four equal annual installments of $5.00 each. Without interest, the $20.00 would be exhausted after making the last payment. If the $20.00 is deposited in the bank at 5% interest, it will be worth $21.00 at the end of the first year. Paying out $5.00 would leave $16.00 to be held at 5% interest for the second year. The $16.00 invested at 5% will earn $0.80 in one year. Subtracting $5.00 from $16.80 leaves $11.80 to be held at 5% interest for the third year. The $11.80 will earn $0.59 by the end of the third year at 5%. Another annual payment of $5.00 leaves $7.39 to earn interest during the fourth year. Adding the $0.37 interest earned during the fourth year and subtracting the final $5.00 annual payment leaves a balance of accrued interest of $2.76. The earning power of money permits another strategy. If the initial deposit were reduced to $17.73 at 5% interest, $5.00 could be paid out each year for four years and nothing would be left. This example illustrates the distinction between the terms “equivalent” and “equal.” The $20.00 is equal to four payments of $5.00 each. It also would be equivalent to four $5.00 annual payments (only) if the interest rate were zero. The $17.75 is not equal to the sum

of four payments of $5.00 each. However, when $17.73 is invested at 5%, it is equivalent to four annual payments of $5.00. The term “equivalent” implies that the concept of the time value of money is applied at some specific interest rate. Therefore, for an amount of money to have a precise meaning, it must be fixed both in time and amount. Mathematical formulas and tables are available to translate an amount of money at any particular time into an equivalent amount at another date. Many kinds of translations are possible. For example, a single amount of money can be translated into an equivalent amount at either a later or an earlier date. This is accomplished by calculating the present worth (PW) or the future worth (FW) as of the present date. Single amounts of money can be translated into equivalent annuities (A) involving a series of uniform amounts occurring each year. Conversely, annuities can be translated into equivalent single amounts at either an earlier or a later date. The present worth of an annuity (P/A) is the single amount of money equivalent to a future annuity. The single amount equivalent to a past annuity is referred to as the future worth of an annuity (F/A). It is also possible to calculate the amount of money that would be equivalent to a nonuniform series of cash flows. Two types of nonlinear series that find application are an arithmetic progression, in which the series changes by a constant amount, and a geometric progression, in which the series changes by a constant rate. The arithmetic progression is considered to be representative of variable costs, such as maintenance costs, which may increase as equipment ages. The geometric progression is used to represent the effects of inflation or deflation.

C. NOTATION AND TERMINOLOGY The American National Standards Institute (ANSI) standard Z94.5 titled “Engineering Economy” consists of a compilation of the symbology and terminology of the field so that the improved communication benefits of standardization are available to practitioners [1]. With the development and publication of this standard and its adoption by the Institute of Industrial Engineers and the Engineering Economy Division of the American Society of Engineering Education, it is expected that future books and articles will utilize these symbols common to engineering economy because they represent the consensus choice of the prominent modern authors and educators in this field. This would avoid one of the significant previous deterrents to the use of these methods in the past. The reader is referred to ANSI standard Z94.5 for further details, including functional forms and uses of compound interest factors, and formulas involving annual compounding and others involving continuous compounding [1].

METHODS OF ECONOMIC ANALYSIS

The basic form of the notation used for all time value factors consists of a ratio of two letters representing two amounts of money (e.g., P/A or F/P) plus an interest rate (i%) and a number of periods (n). The customary manner of writing these is ðP=A; i%; nÞ or ðF=P; i%; nÞ The present worth (P) of a known annuity (A) can also be expressed as follows: P ¼ AðP=AÞ

ð3:1Þ

or the future worth (F) of a known present amount (P) is F ¼ PðF=PÞ

ð3:2Þ

Other forms also are common, such as PW() or FW(), which are called “operators” because they represent some computational operation, such as the PW or the FW of whatever is inside the parentheses. It should be evident that (F/P) is the reciprocal of (P/F) and that (F/A) and (A/F), and so on, also are reciprocals. This observation is useful because it means that only three time value factors need to be tabulated in order to conduct six operations. Another algebraic relationship shows that if two time value factors are multiplied, the product is a third time value factor. For example, ðA=FÞ ðP=AÞ ¼ ðP=FÞ

ð3:3Þ

D. METHODS OF ECONOMIC ANALYSIS Economic analysis methods that are concerned with the entire service life sometimes are called “life-cycle cost” methods. Those that lead to single-measure numbers include: (a) Internal rate of return (IROR) (b) Discounted payback (DPB) (c) Present worth (PW) method, also referred to as the net present value (NPV) (d) Present worth of future revenue requirements (PWRR) (e) Benefit–cost (BCR) ratios All five methods employ the concept of PW. While each method has certain advantages, the individual methods

23

vary considerably with respect to their application and complexity. The IROR method compares the initial capital investment with the PW of a series of net revenues or savings over the anticipated service life. Expenses include all operation, maintenance, taxes, insurance, and overheads but do not include return on (or of) the invested capital. From an economic standpoint, IROR consists essentially of the interest cost on borrowed capital plus any existing (positive or negative) profit margin. The disadvantage of this method is that it ignores benefits extending beyond the assumed life of the equipment and thereby may omit a substantial part of the actual service life. This may lead to unnecessarily pessimistic measures of long-range economy. The PWRR method is particularly applicable to regulated public utilities, at which the rates of return are more or less fixed by regulation. It is particularly applicable when it has already been determined by IROR analysis that a project is economically viable, and the engineer wishes to determine which is the most economic alternative under circumstances wherein several alternatives produce the same revenue but some of them create less expense (requirement for revenue) and consequently a greater profit margin (or lower losses) than others. The principal objection to the PWRR method is that it is inadequate when alternatives are competing for a limited amount of capital because it does not identify the alternative that produces the greatest return on invested capital. The DPB method is somewhat more complicated than the PWRR method. The BCR method is similar to the IROR method because both methods involve assessment of alternatives, not only for economic measures compared with a “do-nothing” scenario, but also for incremental measures associated with incremental capital investments. The PW method, also referred to as the NPV method, is considered the easiest and most direct of the five methods and has the broadest application to engineering economy problems. Many industries refer to this method as the “discounted cash flow” method of analysis. This method often is used as the “referee” method to test the results of other methods of analysis. Under the circumstances, it is not surprising that there is a preference for this method and that primary attention is given to this method in this chapter. Those interested in exploring the other methods are referred to standard texts on the subject. D1. Annual Versus Continuous Compounding Because actual cash flow (both inward and outward) is continuous, it appears that continuous compounding is the more accurate assumption for engineering economy studies. However, although the overall cash flows tend to be continuous, the cash flow data seldom are sufficiently precise in economy studies to take full advantage of continuous

24


compounding. It is also true that the normal purpose of the economy study is to analyze some specific event that will occur at a specific time, so a procedure that is readily associated with a specific time is usually the most appropriate. For these reasons, and because annual compounding is conceptually easier to understand and apply, annual compounding is the method primarily used in this chapter. The PW method is a form of discounted cash flow (DCF), wherein cash flow data, which include dates of receipts and disbursements, are discounted to PW. Before applying these methods, a management judgment is made as to the desired life (usually expressed as a number of years, n) and the minimum acceptable rate of return (ROR) on invested capital for a project (expressed in terms of the effective interest rate, i, or the nominal rate of return, r). Rates of return (before taxes) vary among industries, ranging from 10–15%, where obsolescence is not high, to 25–40% (or perhaps higher) for dynamic industries. Obviously, it is convenient if the minimum acceptable ROR is known before making engineering economy analyses. However, it is not always easy to learn what is considered an acceptable ROR. Under such circumstances, it has been helpful to prepare a series of economic alternatives in which the ROR is varied so that management can make a choice. An example illustrates the features of the method. In this case, the “longhand” method is used. In practice, there are several shortcuts. Should an expenditure of $15,000 be made to reduce labor and maintenance costs from $8200 to $5100 per year? Money is worth 10% and the life of the project is 10 years. For simplicity, the effects of taxes and depreciation have been neglected in this example. Table 3.1 shows the projected pattern of cash flow for the “defender” (the present method) and the “challenger” TABLE 3.1. Tabulation of Cash Flow Period (year)

Plan A (dollars) (Defender)

Plan B (dollars) (Challenger)

B–A (dollars)

0 1 2 3 4 5 6 7 8 9 10 Totals

— 8,200 8,200 8,200 8,200 8,200 8,200 8,200 8,200 8,200 8,200 82,000

15,000 5,100 5,100 5,100 5,100 5,100 5,100 5,100 5,100 5,100 5,100 66,000

15,000 þ 3,100 þ 3,100 þ 3,100 þ 3,100 þ 3,100 þ 3,100 þ 3,100 þ 3,100 þ 3,100 þ 3,100 þ 16,000

Source: E. Verink, “Corrosion Economic Calculations,” ASM Handbook, Vol. 13, 9th ed., Corrosion, ASM International, Materials Park, OH, 1987, p. 372 (Table 4).

(the proposed method) over the life of the project. A minus sign means that money leaves the “bank,” whereas a plus sign means that the “bank balance” increases in size. The net cash flow for each year appears in the right-hand column under the heading “B A.” It is apparent that selection of the challenger (plan B) results in a net positive cash flow. That is, the net amount of money in the bank is increased over the life of the project when plan B is selected. Before reaching a conclusion regarding implementation of plan B, these cash flows are reduced to a common basis for comparison and to determine whether the objective of a 10% ROR has been achieved. The PW (or present value) of these cash flows provides such a basis. The ROR for plan B is calculated by iteration using interest tables and interpolating between values. D1.1. First Iteration. Assume a rate of return of 10% and refer to Table 3.2: ðPWÞ ¼ 15; 500 þ 3100ðP=A; 10%; 10 yearsÞ

ð3:4Þ

¼ 15; 500 þ 3100ð6:145Þ ¼ þ $4047:95 ð3:5Þ The ROR for which the discounted cash flow is equal to zero (i.e., the first term on the right of the above expression is balanced by the second term) is the actual rate of return. From the first iteration, it already is evident that plan B returns more than 10% because the net cash flow is positive. Thus, plan B, the challenger, is the more economical. The numerical value of the actual rate of return can be determined by additional iterations. Such an exercise reveals that the ROR in this case is 16.1%. In this example, plans A and B could represent alternative materials of construction having different corrosion rates, with the annual dollar difference being related to the consequences of corrosion on maintenance costs. This presupposes the availability of corrosion data that can be used to estimate expected life, maintenance costs, and so on. Other examples illustrate how to account for the effects of salvage value, taxes, and depreciation.

E. DEPRECIATION Depreciation has been defined as the lessening in value of an asset with the passage of time. All physical assets (with the possible exception of land) depreciate with time. There are several types of depreciation. Two of the more common types are physical depreciation and functional depreciation. Accidents can also cause loss of value, but this cause is often accommodated in other ways (i.e., insurance or reserves) and is not considered here. Physical depreciation includes such phenomena as deterioration resulting from corrosion, rotting of wood, bacterial

DEPRECIATION

TABLE 3.2.

25

10% Interest Factors for Annual Compounding Single Payment

Equal-Payment Series

CompoundAmount Factor (To find F Given P F/P, i, n)

PresentWorth Factor (To find P Given F P/F, i, n)

CompoundAmount Factor (To find F Given A F/A, i, n)

Sinking-Fund Factor (To find A Given F A/F, i, n)

Present-Worth Factor (To find P Given A P/A, i, n)

Capital Recovery Factor (To find A Given P A/G, i, n)

Uniform Gradient-Series Factor (To find A Given G A/G, i, n)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 40 45 50 55 60 65 70 75 80 85

1.100 1.210 1.331 1.464 1.611 1.772 1.949 2.144 2.358 2.594 2.853 3.138 3.452 3.798 4.177 4.595 5.054 5.560 6.116 6.728 7.400 8.140 8.954 9.850 10.835 11.918 13.110 14.421 15.863 17.449 19.194 21.114 23.225 25.548 28.102 45.259 72.890 117.391 189.059 304.482 490.371 789.747 1,271.695 2,048.400 3,298.969

0.9091 0.8265 0.7513 0.6830 0.6209 0.5645 0.5132 0.4665 0.4241 0.3856 0.3505 0.3186 0.2897 0.2633 0.2394 0.2176 0.1979 0.1799 0.1635 0.1487 0.1351 0.1229 0.1117 0.1015 0.0923 0.0839 0.0763 0.0694 0.0630 0.0573 0.0521 0.0474 0.0431 0.0392 0.0356 0.0221 0.0137 0.0085 0.0053 0.0033 0.0020 0.0013 0.0008 0.0005 0.0003

1.000 2.100 3.310 4.641 6.105 7.716 9.487 11.436 13.579 15.937 18.531 21.384 24.523 27.975 31.772 35.950 40.545 45.599 51.159 57.275 64.003 71.403 79.543 86.497 96.347 109.182 121.100 134.210 148.631 164.494 181.943 201.138 222.252 245.477 271.024 442.593 718.905 1,163.909 1,880.591 3,034.816 4,893.707 7,887.470 12,708.954 20,474.002 32,979.690

1.0000 0.4762 0.3021 0.2155 0.1638 0.1296 0.1054 0.0875 0.0738 0.0628 0.0540 0.0468 0.0408 0.0358 0.0315 0.0278 0.0247 0.0219 0.0196 0.0175 0.0156 0.0140 0.0126 0.0113 0.0102 0.0092 0.0083 0.0075 0.0067 0.0061 0.0055 0.0050 0.0045 0.0041 0.0037 0.0023 0.0014 0.0009 0.0005 0.0003 0.0002 0.0001 0.0001 0.0001 0.0000

0.9091 1.7355 2.4869 3.1699 3.7908 4.3553 4.8684 5.3349 5.7590 6.1446 6.4951 6.8137 7.1034 7.3667 7.6061 7.8237 8.0216 8.2014 8.3649 8.5136 8.6487 8.7716 8.8832 8.9848 9.0771 9.1610 9.2372 9.3066 9.3696 9.4269 9.4790 9.5264 9.5694 9.6086 9.6442 9.7791 9.6628 9.9148 9.9471 9.9672 9.9796 9.9873 9.9921 9.9951 9.9970

1.1000 0.5762 0.4021 0.3155 0.2638 0.2296 0.2054 0.1875 0.1737 0.1628 0.1540 0.1468 0.1408 0.1358 0.1315 0.1278 0.1247 0.1219 0.1196 0.1175 0.1156 0.1140 0.1126 0.1113 0.1102 0.1092 0.1083 0.1075 0.1067 0.1061 0.1055 0.1050 0.1045 0.1041 0.1037 0.1023 0.1014 0.1009 0.1005 0.1003 0.1002 0.1001 0.1001 0.1001 0.1000

0.0000 0.4762 0.9366 1.3812 1.8101 2.2236 2.6216 3.0045 3.3724 3.7255 4.0641 4.3884 4.6988 4.9955 5.2789 5.5493 5.8071 6.0526 6.2861 6.5081 6.7189 6.9189 7.1085 7.2881 7.4580 7.6187 7.7704 7.9137 8.0489 8.1762 8.2962 8.4091 8.5152 8.6149 8.7086 9.0962 9.3741 9.5704 9.7075 9.8023 9.8672 9.9113 9.9410 9.9609 9.9742

90 95 100

5,313.023 8,556.676 13,780.612

0.0002 0.0001 0.0001

53,120.226 85,556.760 137,796.123

0.0000 0.0000 0.0000

9.9981 9.9988 9.9993

0.1000 0.1000 0.1000

9.9831 9.9889 9.9928

n

Source: G. J. Thuesen and W. J. Fabrycky, Engineering Economy, 6th ed., Prentice-Hall, Englewood Cliffs, NJ, 1984, p. 574.

26


action, chemical decomposition, wear and tear, and so on, which can reduce the ability of an asset to render its intended service. Functional depreciation results not from the inability of an asset to be available to serve its intended purpose, but rather from the fact that some other asset is available that can perform the desired function more economically. Thus, obsolescence and/or inadequacy, or inability to meet the demands placed on the asset, lead to functional depreciation. Technological advances produce improvements that often result in obsolescence of existing assets. The manner in which depreciation is accounted for is largely a tax question. The language of the tax laws specifies the procedures that are permissible. Tax laws change from time to time, so procedures that are attractive under a given set of circumstances may become unattractive (or even forbidden) under other circumstances. Some of the more common methods of depreciation include: 1. 2. 3. 4. 5.

Straight line Declining balance Declining balance switching to straight line Sum of the year’s digits Accelerated cost recovery system

E1. Straight-Line Method The straight-line depreciation method assumes that the value of an asset declines at a constant rate. If the asset originally cost $6000 and had a salvage value of $1000 after five years’ life, the annual depreciation would be ($6000 – 1000)/5 ¼ $1000/year. Symbolically, this can be expressed Annual depreciation;

D¼

PS n

ð3:6Þ

The “book value” of the asset would decrease at the end of each year by the amount of the annual depreciation until, at the end of the fifth year, the book value would be the same as the salvage value. E2. Declining-Balance Method The declining-balance method assumes that the asset depreciates more rapidly during early years than in later years. A certain percent of depreciation is applied each year to the remaining book value of the asset. Under these circumstances, the size of the depreciation declines each successive year until the asset is fully depreciated. When the declining-balance method is used, the maximum rate that has been permissible for tax purposes is double the straightline rate. This accounts for the term “double-declining balance.”

E3. Declining Balance Switching to Straight Line Prior to 1981 in the United States, it was allowable to depreciate an asset using declining-balance depreciation for the early years and then switch to straight-line depreciation when the allowable depreciation (using declining balance) falls below the amount permissible under straight-line depreciation. Switching to straight-line depreciation permits the book value to go to zero eventually. The decliningbalance method switching to straight line has been incorporated in part of the 1986 version of the ACRS depreciation method discussed below. E4. Sum-of-the-Year’s-Digits Method The sum-of-the-year’s-digits method assumes that the value of an asset decreases at a decreasing rate. Assume an asset has a five-year life. The sum of the year’s digits equals 1 þ 2 þ 3 þ 4 þ 5 ¼ 15. For the $6000 asset mentioned above, which has a salvage value after five years of $1000, the first year’s depreciation would be ($6000 1000)(5/15) ¼ $1666.67. The second year’s depreciation would be ($6000 1000) (4/15) ¼ $1333.33, and so on, until the asset is fully depreciated. This method is no longer permitted in the United States for new assets. Most of the property acquired before 1981 is still being depreciated by one of the methods mentioned above. Each of these (pre-1981) methods involves the taxpayer estimating a “useful life” either on the basis of experience or based on a guideline from the U.S. Internal Revenue Service in Washington, DC. The guidelines are presented as ranges of values in the class life asset depreciation range (ADR) system. The midpoint of the range is referred to as the “ADR life.” E5. Accelerated Cost Recovery System Method For property placed in service after 1981, the accelerated cost recovery system (ACRS), which is a part of the Economic Recovery Act of 1981, prescribes a different method for recovering the cost of depreciable property. The Tax Reform Act of 1986 revamped the ACRS enacted in 1981. The 1986 act generally is less generous than the prior law, which combined ACRS with investment tax credit (ITC). While ACRS was originally designed primarily as an incentive for investment, the new rules are intended to provide a more even match between the class lives of particular assets and their useful lives. Nonetheless, deductions are more accelerated than under pre-1981 law. The principal differences between the prior version of ACRS and the new law are in the class lives of assets and the methods of recovering their costs. As in the prior law, salvage continues to be disregarded. The new system is generally effective for assets placed in service after 1986, but by special election, property placed in service after July 1, 1986, also may qualify under the new rules.

DEPRECIATION

TABLE 3.3.

27

Recovery Property

ACRS Class and Method

ADR Midpoint

Special Rules

3-year, 200% DB 5-year, 200% DB

4 years or less > 4 to Ep, pitting will occur; E < Erep, growing pits will repassivate.

More detailed reviews of the various electrochemical methods used to measure the susceptibility to pitting corrosion are given by Szlarska-Smialowska [1] and Sedriks [2]. Obviously, these quantities mentioned above can be measured as functions of the alloy composition or of the composition of the environment. Without discrediting the practical value of such investigations, especially in comparing alloys as well as environments with respect to their pittingsusceptibility or pitting-promoting tendency, they give neither any direct insight into the mechanisms of localized corrosion processes nor can the values be used as true limiting potential values to prevent localized corrosion processes in engineering applications. B2. Effects of Alloying and Microstructure The alloy composition as well as the microstructure can have a strong influence on the pitting resistance of an alloy, as shown, for instance, by Horvath and Uhlig [3], who demonstrated the beneficial effect of chromium and molybdenum in stainless steels. The pitting potential was found to increase dramatically with chromium contents > 20 wt %, whereas molybdenum is effective at minor concentrations of 2–6 wt %, but only in the presence of chromium. Similar effects have also been reported for small amounts of alloyed nitrogen or tungsten. Various explanations have been given to explain the strong influence of molybdenum on the pitting behavior of stainless steels, which is also well confirmed in engineering practice. It was suggested that molybdenum is adsorbed on the surface as molybdate or acts by blocking active surface sites, inhibiting active metal dissolution and finally favoring repassivation [4–6]. Other models suggest that molybdenum, as well as other elements, improves the cation-selective properties of the passive film, hindering the migration of aggressive anions, such as chlorides, to the metal surface [7, 8] or reduce the flux of cation vacancies in the passive film [9]. Though the exact mechanism is not clear yet, most of the research done so far favors an effect on the growth stage rather than on the initiation stage of localized corrosion. Recent results, applying microelectrochemical techniques, confirmed that even in the superaustenitic stainless steels molybdenum strongly improves the repassivation behavior but has no influence on pit initiation [10, 11]. Similar effects can also be observed on aluminum, where small additions of alloying elements may increase the pitting potential as long as the structure is single phase. Aluminum–copper alloys, widely used in the aircraft industry, are well known and intensively studied examples with respect to their pitting behavior [12]. The pitting potential increases with increasing copper concentration as long as copper is in solid solution. Since aluminum usually exhibits rather low solubility limits, alloying elements tend to form second-phase intermetallic compounds, such as Al2Cu. In the presence of

STAGES OF LOCALIZED CORROSION

Al2Cu the pitting resistance decreases to values of aluminum with small copper concentrations. The corresponding decrease in pitting potential is explained by the existence of a copper-depleted zone around these particles, where pits would initiate first. Localized corrosion of passive metals almost always initiates at local heterogeneities, such as inclusions and second-phase precipitates as well as grain boundaries, dislocations, flaws, or sites of mechanical damage. In the case of stainless steel surfaces, pit initiation occurs almost exclusively at sites of MnS inclusions, which are found in commercial as well as in high-purity alloys. A more detailed discussion of this subject will follow. To prevent inclusions and precipitates, nonequilibrium single-phase conditions can be attained by special preparation techniques, such as rapidly quenching or physical vapor deposition. The resulting microstructure is either nanocrystalline or amorphous. It was recently shown that sputter-deposited aluminum alloys containing only a few atomic percent of metal solute such as Cr, Ta, Nb, W, Mo, or Ti exhibit a strong increase of Ep of 0.2–1V [13–17]. Similar results were also obtained with sputter-deposited stainless steels, where nonequilibrium single-phase structures with molybdenum concentrations up to 14 wt % could be obtained [18, 19]. In both cases, the increase in pitting resistance was explained by the reduced pit initiation tendency as well as by a more protective passive film, favoring rapid repassivation. B3. Effect of Temperature Increasing temperature usually also increases the pitting tendency of metals and alloys. At low temperature, high pitting potentials are observed. In the case of stainless steels, a strong decrease of pitting potentials in the temperature range between 0 and 70 C of 0.5 V can be observed as shown in Figure 13.2 [20, 21]. This strong dependence on temperature has led to experimental techniques which allow stainless steels to be ranked according to their pitting susceptibility. A critical pitting temperature (CPT) has been defined, below which a steel in an aggressive Cl containing solution, usually a FeCl3 solution, would not pit regardless of potential and exposure time [20, 22]. Furthermore, the CPT used to characterize the pitting resistance of alloys is often correlated with the composition of stainless steels, especially with Cr, Mo, and N [1, 2, 20]. The effect of temperature on the pitting behavior of other metals and alloys such as Ni, Al, Ti, and its alloys has also been studied but to a lesser extent. Important earlier results are summarized in [20]. Nevertheless most of the investigations on temperature effects of localized corrosion deal with the significance of temperature to evaluate and compare pitting susceptibilities of different metals and alloys rather than with the more fundamental aspects of temperature effects on pit initiation and pit growth mechanisms.

159

FIGURE 13.2. Effect of temperature on pitting potential Ep of 304 stainless steel (SS) in 0.1M NaCl with different sulfur contents [21].

B4. Stochastic Aspects Corrosion processes are based mainly upon deterministic approaches, such as the electrochemical theory of corrosion. Localized corrosion events, however, due to their unpredictable occurrence, cannot be explained without using statistical methods to evaluate the experimental data. Stochastic aspects of pitting corrosion were studied in the late 1970s, especially by authors in Japan [23–26]. Recently, a convincing review of statistical and stochastic approaches to localized corrosion was published by Shibata [27]. He evaluated large numbers of pitting potential values using a Gaussian distribution, whereas the Poisson distribution was found to be a better approach for pit generation. The results indicate that different pit generation rates can be observed as a function of time. He proposed two groups of models considering either pit generation events alone or assuming pit generation and subsequent repassivation processes. The latter model could be fitted more satisfactorily to the various cases studied experimentally. More detailed information will be given in Chapter 27.

C. STAGES OF LOCALIZED CORROSION The different stages of localized corrosion can best be explained and discussed in connection with its potential dependence. In Figure 13.3, a typical potential current curve of a passive metal, such as stainless steel, measured in chloride solutions, shows the different stages of pitting corrosion. At lower potentials, pit initiation is followed by rapid repassivation. This stage is usually referred to as metastable pitting. The resulting current transients differ widely with

160

LOCALIZED CORROSION OF PASSIVE METALS

FIGURE 13.3. Typical potential current curve of stainless steel in chloride solution showing different stages of localized corrosion (see text).

respect to the peak current height as well as the lifetime. Small current transients in the femtoampere (fA) and picoampere (pA) range, corresponding to pits of nanometer and micrometer size can be detected only by applying micro- and nanoelectrochemical methods, as will be discussed later. Increasing the potential generally leads to larger current transients with higher peak currents and longer lifetimes, indicating an extended pit growth period. Above a certain potential or potential range, a transition to stable pit growth occurs. Even above the pitting potential, repassivation may still occur, showing the stochastic character of localized corrosion processes. Whether well-defined pitting potentials can be determined depends on a variety of factors, such as the type of metal, the chloride concentration, and the temperature. The higher the chloride concentration or the temperature, the more precise the resulting pitting potentials usually are. Furthermore the use of crevice-free experimental techniques may also be very decisive in obtaining reproducible results. Metastable pitting is usually not considered as a real corrosion risk from an engineering point of view. Nevertheless, studies of metastable pitting as a precursor to stable pitting may provide valuable insights into fundamental aspects of pitting corrosion, since pit initiation as well as the trans formation of metastable into stable pits are key factors

in localized corrosion processes. Investigations of metastable pitting also allow a statistical evaluation of corrosion data necessary to study stochastic pitting models. With respect to engineering application, studies of metastable pitting may also substantially improve the evaluation of metal–environment systems. The occurrence of metastable pitting below the pitting potential indicates a potential corrosion risk, especially when crevice conditions cannot be completely excluded. Furthermore the repassivation behavior of metals and alloys can easily be studied, which is very important for developing highly corrosion-resistant alloys.

D. METASTABLE PITTING: PIT INITIATION AND REPASSIVATION A number of models have been proposed to describe the initiation of localized corrosion of passive metals based either on the breakdown processes of the passive film itself or on structural defects or heterogeneities of the underlying metal or alloy such as dislocations, grain boundaries, secondphase precipitates, or nonmetallic inclusions. Certainly a strict differentiation of the two approaches is not always possible, since film breakdown and structural parameters of the underlying metal may be correlated.

METASTABLE PITTING: PIT INITIATION AND REPASSIVATION

161

FIGURE 13.4. Models for pit initiation leading to passive film breakdown. [After Strehblow and coworkers [28–30]. Reproduced with permission from H. Kaesche, Die Korrosion der Metalle, 3rd edition. Springer, Berlin, Germany, 1990, Fig. 12.12a–c, p. 308. Copyright Springer.]

D1. Passive Film Breakdown Theoretical models that describe the initiation process leading to passive film breakdown may be grouped into three classes: (1) adsorption and adsorption-induced mechanisms, where the adsorption of aggressive ions like Cl is of major importance, (2) ion migration and penetration models, and (3) mechanical film breakdown theories, as shown in Figure 13.4 [28–30]. In the case of the cluster adsorption model, originally proposed by Heusler and Fischer [31] several years ago for iron, localized adsorption of chlorides leads to an enhanced oxide dissolution at these sites with subsequent thinning of the oxide film until finally a complete removal is achieved and active dissolution starts. During the latter years further refinements of this model were obtained evaluating the measured induction times statistically [32]. Further evidence for the stochastic nature of these initiation processes were also obtained on stainless steels [10]. For the mechanical breakdown of passive films models have been discussed, where the breakdown is either the principal step, giving the environment direct access to the metal surface, as shown by Sato [33], or an additional step combined with other processes as in the case of the defect model as pointed out by Macdonald and co-workers [34, 35]. A critical review of these mechanisms is given in [36]. Most of the theoretical models for pit initiation have not been sufficiently verified experimentally. Since the relationships used to test the different models are very general in nature, such as the correlation between pitting potential and chloride concentration or the quantitative prediction of induction times, they do not really cover specific model-

sensitive aspects. Furthermore, the pitting potentials proposed in these models are actually critical pit initiation potentials, which often do not correspond with the experimentally determined pitting potentials, at which the transition to stable pit growth takes place. Pit initiation processes, on the other hand, often occur at much lower potentials, as was recently shown on stainless steels by applying a new microcapillary technique to measure current transients in the picoampere and femtoampere range [10, 37, 38]. Recently, new results on the initiation and formation of porous semiconductors were obtained [39]. In case of n- and p-type GaAs, local dissolution can be electrochemically triggered in chloride-containing solutions resulting in a porous structure. In contrast to pitting corrosion of passive metals, local dissolution can be achieved not only by the presence of an oxide film (p-type GaAs) but also by depletion conditions in the semiconductor space-charge layer (n-type GaAs). In the latter case, the presence or absence of an oxide film is not significant for the pitting process. This is of particular significance since the previously discussed theoretical approaches ascribe a key role in the localized nature of pitting to the properties of a surface oxide film.

D2. Structural Parameters Numerous investigations during recent years have shown that the sites of pit initiation on passive metal surfaces may generally be related to defect structures of the underlying metals. Detailed summaries are given in [1, 2], In case of Ni single crystals, for instance, it has been clearly demonstrated that the emerging points of screw dislocations are especially

162


susceptible sites for pit nucleation [40]. Furthermore, the shape as well as the pit density depend markedly on the crystallographic orientation of the surface exposed. Pits in aluminum, titanium, and their alloys are typically associated with intermetallic phases in these alloys. Aluminum contaminated with iron exhibits an increased pitting susceptibility due to the presence of FeAl3 particles that act as local cathodes on the metal surface [41–43]. Recent investigations using microsensors to measure local pH distributions revealed a buildup of hydroxide ions over FeAl3 due to the cathodic reaction on the particle surface. This may lead to alkaline dissolution of the matrix at the particle interface [44]. Preferential attack can also occur on the surface of intermetallic phases, such as on Mg2Al3 and MgSi in Al–Mg and Al–Mg–Si alloys [45, 46], In Ti–Al alloys it is assumed that particles of Ti3Al2 are effective in initiating pitting corrosion [47]. On the other hand, nonmetallic inclusions may also act as potential nucleation sites [48–50]. Sulfide inclusions in stainless steels are particularly susceptible. Figure 13.5 shows the initial stage of localized corrosion on stainless steel at the site of a MnS inclusion [11]. This is observed not only in austenitic [51–53] but also in ferritic [54, 55] stainless steels. These inclusions are often manganese sulfides or manganesecontaining sulfide compounds [53–56]. It has been suggested that, at low manganese levels in steel, CrS is the thermodynamically stable sulfide, while above some level of manganese the stable sulfide is an iron–manganese spinel, which appears to be a better initiation site for pitting than CrS. In a recent survey, Srivastava and Ives [57] summarized the different types of attack on nonmetallic inclusions in stainless steels from a phenomenological point of view.

FIGURE 13.5. Pit initiation and early pit growth at MnS inclusion in 304 SS [11].

Powerful electrochemical techniques have been developed for studying localized phenomena on passive metal surfaces, such as the scanning methods, extensively applied and discussed by Isaacs and co-worker [58, 59]. To improve the local resolution substantially, small area measurements are usually carried out, using either the photoresist techniques for masking off small areas or embedded wires with small cross sections (diameter 50 mm) [60]. Recently, a new microelectrochemical technique applying microcapillaries as electrochemical cells has been developed by Suter and coworkers [10, 11, 37, 38]. Only small surface areas with a few micrometers or even nanometers in diameter are exposed to the electrolyte. This leads to a strongly enhanced current resolution, down to picoamperes and femtoamperes. Therefore local processes in the micrometer and nanometer range can easily be studied. In the case of stainless steels, the pit initiation process due to the oxidative dissolution of active MnS inclusions as a precursor of pitting corrosion can be investigated directly as shown in Figure 13.6. Additionally, the results clearly indicate that the dissolution of inclusions takes place even in chloride-free solutions,

FIGURE 13.6. Local potentiodynamic polarization curves at active and at inactive MnS inclusion site with corresponding scanning electron microscopy (SEM) photographs taken after polarization measurements [38].

METASTABLE PITTING: PIT INITIATION AND REPASSIVATION

163

FIGURE 13.7. Correlation between pitting potential and inclusion size [10].

whereas chlorides are required for metal dissolution and stable pit growth processes. The results in Figure 13.6 show that the microcapillary technique also makes it possible to distinguish between active and inactive inclusions with respect to pit initiation. An increase of the sulfur content drastically increases the number of large current transients due to an increasing number of large MnS inclusions, verified by microstructural investigations [37, 38]. Similar effects have also been observed on iron in contact with Cl -containing borate solutions [61]. Noise analysis obtained from microelectrochemical investigations of stainless steels under potentiostatic conditions revealed that the current noise, expressed as standard deviation si of the passive current, increases linearly with the size of the exposed area, whereas the pitting potential decreases. Computer simulations showed [10, 38] that the current noise, largely caused by the dissolution of small inclusions, can be correlated to the size of the inclusions. Therefore the pitting potential is also related to the size of active inclusions [62], as shown in Figure 13.7. Specifically the size of the inclusions in stainless steels has to be kept below 1 mm to improve substantially the pitting resistance of stainless steel. This effect was already observed much earlier by simple immersion tests [63]. The effect of molybdenum on the pitting behavior of stainless steels can also be studied by microelectrochemical techniques and then compared to molybdenum free alloys having approximately the same impurity level, as shown in Figure 13.8 [10]. Molybdenum in superaustenitic stainless steels has only a minor effect on the initiation process. The superior corrosion resistance of these high-molybdenum-

FIGURE 13.8. Microtransients of DIN 1.4301 (0% Mo) and DIN 1.4529 (6.4% Mo) stainless steels in chloride-free and chloridecontaining solutions [10].

containing alloys rather has to be attributed to a considerably improved repassivation behavior. D3. Stability of Passive Films The stability of passive films plays an important role with respect to the pitting behavior of passive alloys. Fast and effective repassivation of locally activated metal surfaces only occurs if a stable passive film is formed. Since pit initiation can hardly be neglected on commercial alloys in real environments, rapid repassivation in which a passive film of high stability is formed, is very important to obtain highly corrosion-resistant alloys. In order to understand the importance of the chemical and electrochemical stability of passive films with respect to localized corrosion, the behavior of naturally grown as well as synthetically prepared thin oxide films has to be studied [64, 65]. In several contributions Schmuki and co-workers [66, 67] showed that the semiconductive properties and the chemical stability of the thin oxide films of iron and chromium correlate quite well with those of naturally grown films. The results indicate that the presence of Fe(II), acting as doping species in iron oxide, strongly effects its stability. The good correlation between the semiconductive properties and the stability of passive films was

164


E1. Growth Kinetics

FIGURE 13.9 Dissolution rate of sputter-deposited Fe/Cr oxide films as function of Cr2O3 content during galvanostatic reduction and oxidation in borate buffer, pH 8.4. Dissolution rates determined from the drop of the edge height of XANES spectra during the experiment. (Reproduced by permission of The Electrochemical Society, Inc. [70]).

additionally confirmed, using in situ X-ray absorption nearedge spectroscopy (XANES) techniques [68, 69]. The experiments showed that the mixed oxides are far more stable than the pure oxides as shown in Figure 13.9. The presence of sufficient amounts of chromium oxide protects iron oxide against reductive dissolution, whereas iron oxide protects chromium oxide from oxidative dissolution. During anodic polarization of iron oxide films in acidic solutions, a deleterious effect of chloride anions compared with sulfates was found. In HCl solutions of increasing concentration, not only the increased acidity but also the increased chloride concentration accelerates the dissolution markedly [70]. E. PIT GROWTH Whentheactivestatewithinpitsandcrevicesismaintainedover an extended period of time, rapid metal dissolution usually occurs. The resulting pit and crevice geometries as well as the surface state within the pits vary markedly from open and polished hemispherical pits on free surfaces to etched cracklike shapes within crevices, depending largely on the type of rate-controlling reactions during the growth stage. The kinetics of pit growth is not only of scientific interest but is also of great engineering importance for commercial corrosion-resistant materials, since the possibility of local breakdown should not be ignored under practical conditions. Sufficient knowledge of the mechanisms of growth and stability is therefore a necessary requirement in order to predict the corrosion behavior of passive metals correctly as well as for developing new corrosion-resistant materials.

Although pit growth is experimentally much easier to study and quantify than pit nucleation, substantive information on pit growth mechanisms was mainly obtained during the last two decades. In earlier research on aluminum [71–79], iron [80], and stainless steels [81, 82] under open-circuit conditions, studies were limited to the description of growth rate by a simple power law dp ¼ atb, whereby dp is the pit depth, t is time, and a and b are constants, with the latter averaging in many cases 0.5. Values of b < 1 clearly indicate a decreasing growth rate with time as shown, for example, on aluminum in various tap waters [79]. A comprehensive literature survey, summarizing the rate laws of various metals and alloys is given in [83]. Numerous investigations have confirmed that the presence of chlorides is necessary for pit growth and that the growth rate increases with increasing chloride concentration [84–90]; other anions, however, behave differently. To overcome the problems associated with accurate pit current density measurements, special techniques have been developed to determine pit growth kinetics. Hunkeler and coworkers [83, 91, 92] used the time for pits to penetrate metal foils of different thickness to determine the growth rate of the fastest growing pits. Using this simple method pit growth rates in aluminum as well as stainless steels were measured and several parameters, such as the potential, the composition of the electrolyte, and the temperature, were varied. Another elegant approach to study the growth kinetics involved the investigations of two-dimensional pits in thin metal films as demonstrated by Frankel et at. [93–96]. The measurements of lateral pit growth rates from analysis of images of the growing two-dimensional pits provides a very simple and direct way via Faraday’s law, with no need for any further assumptions. Investigations were carried out on different materials such as nickel and aluminum alloys. The study of single pits, formed in different ways have also been performed either by masking off a small area, implanting an activating species at a small spot, or using single-pit electrodes, such as embedded wires [60, 97–99]. E1.1. Diffusion Control. Several detailed studies on the kinetics of growing pits, performed in the 1970s, recognized the presence as well as the importance of salt layers during pit growth [100–104]. In Figure 13.10, the pit and crevice growth of 304 SS in a chloride-containing solution as a function of the potential, using the foil technique as mentioned above, is shown [105, 106]. At high potentials, where pits grow on open surfaces as well as crevices, a potential independent growth rate is observed, suggesting diffusion-controlled growth mechanism. When lower potentials are applied, the mode of localized corrosion changes to etch-type crevice corrosion with a strong potential dependence. Below a critical growth rate,

PIT GROWTH

165

FIGURE 13.10. Mean pit and crevice growth rate vp vs. potential of type 304 SS [105, 106, 112]. (Reproduced with permission from Advances in Localized Corrosion, p. 70, Copyright NACE International. All rights reserved.)

depending on the chloride concentration, etch-type crevice corrosion is the only stable form of localized corrosion, since saturation conditions and salt film formation are not attained. Concerning the practical significance of diffusion-controlled pit growth, it must be pointed out that this type of pit growth is not often observed on stainless steels since it usually requires high potentials and/or high chloride concentration. In the case of unidirectional, single-pit growth using wire electrodes, the diffusion-controlled pit growth can easily be calculated using Fick’s first law. Excellent experimental proof was obtained on stainless steel and nickel, for which a parabolic rate law was observed [107]. Additionally, it was found that if the bulk electrolyte contains the corresponding diffusing metal ion in solution, the pit growth is much faster in less concentrated solutions, as expected from theoretical considerations [108]. E1.2. Ohmic/Charge Transfer Control. As pointed out by Beck [109], pitting on titanium and aluminum occurs at high ohmic limited current densities. Generation of large amounts of hydrogen bubbles within the pit strongly increases the mass transport rate. Therefore the fluid flow of the bulk electrolyte has little effect on pit growth under such conditions [110]. Further and more detailed support for

ohmic-controlled pit growth on aluminum was obtained by Hunkeler [83]. For small Tafel constants, as in the case of aluminum, and sufficiently large pits (>10 mm), contributions from charge transfer as well as ohmic transport outside the pit may be neglected, and a simple parabolic rate law [92] can be derived in which the preexponential factor depends directly on the electrolytic conductivity of the bulk electrolyte. Due to the generation of hydrogen bubbles during pitting of aluminum, no significant change in the composition of the electrolyte within the pit takes place, in contrast to situations in which diffusion processes control pit growth. These findings are in excellent agreement with the evaluation of long-term pit growth measurement under open-circuit conditions on aluminum in tap water of known conductivity, as shown in Figure 13.11 [111, 112]. The experimental values fit very nicely into the parabolic rate law. The resulting potential difference of 15 mV is in good agreement with the open-circuit conditions in the absence of strong oxidants. For materials such as stainless steels or nickel, the charge transfer reaction cannot be neglected, as shown in Figure 13.10. A semilogarithmic relationship between current density and potential is observed at low potentials, indicating that mixed ohmic/charge transfer control is probably effective. Similar results obtained by Newman [6] using

166


FIGURE 13.11. Pit growth on aluminum in tap water at opencircuit conditions [111]. (Reproduced with permission. Copyright NACE International. All rights reserved.)

artificial pit electrodes showed that molybdenum in stainless steels ennobles the anodic dissolution reaction in the pit environment.

F. STABILITY OF LOCALIZED CORROSION The concepts of stability of pitting and crevice corrosion are based mainly on either compositional changes of the electrolyte within the pit, salt film formation, or sufficiently large ohmic potential drops within pits. Various authors have suggested that critical concentrations of ionic species must be exceeded for stable pit growth to occur. Galvele [113–115] considered localized acidification as the main reason for stable pit growth. Sato [116, 117] pointed out that two models of metal dissolution have to be distinguished: (1) active or etching dissolution, which occurs at lower potentials, and (2) transpassive or brightening dissolution, which occurs at relatively noble potentials. For etch pitting, a critical hydrogen ion concentration should be reached, whereas for brightening pits, the necessity of a critical concentration of aggressive anions, such as chlorides, was proposed. On the other hand Vetter and Strehblow [118, 119] concluded from theoretical considerations that during the early stage of growth of an open hemispherical pit, the metal chloride concentration within the pit increases, but not sufficiently for precipitation to occur, which they considered as a possible requirement for the stability of pits. Instead, they proposed that an ion-conducting salt layer on the metal surface results from adsorption of aggressive anions, such as chlorides. The significance of an ohmic potential drop within pits was also questioned by the same authors [118, 119]. For iron, they estimated potential drops of 18 mV, which again cannot in any case explain the stability of growing pits. However, several groups at about this time also measured

large potential variations within crevices and pits [120–124], Pickering and Frankenthal [121–123] explained the large potential drops by the presence of high resistance paths caused by gas bubbles inside the pits. More recently, Pickering and co-workers [125, 126] showed that the ohmic (IR) drop, rather than a composition variation, was responsible for stabilizing crevice corrosion in some systems which exhibit active/passive transitions in the bulk solution. Most of these experiments were performed on iron or nickel in sulfuric acid, where the metal at open-circuit conditions is in the active state. Therefore, these results may not be sufficiently general to apply to localized corrosion processes of originally passive metal surfaces, where an activation process has to occur first and the active state has to be stabilized even without an active–passive transition under bulk conditions. The effect of salt films formed within the pit during the initiation or growth process has been discussed frequently [84, 118, 119]. Beck and Alkire [103] and Beck and Chan [127] demonstrated convincingly that the formation of salt films on stainless steels may be important for the stability of growing pits. In using a flow cell, it was shown that with increasing flow velocity the dissolution current increases at first, before at the highest velocity the current drops drastically due to the removal of the salt film and the subsequent repassivation of the surface. Further support for the importance of salt film formation was also obtained from evaluation of metastable pitting on stainless steels [128]. Transition from metastable to stable pit growth occurs only when salt-film precipitation takes place. Additional results showing the significance of a salt film were obtained from electrochemical impedance spectroscopy. The investigation of the anodic dissolution of artificial pit electrodes indicated that, in the case of stainless steel and nickel, significant ohmic potential drops are present within the salt film during diffusion controlled pit growth. The thin, conducting salt film stabilizes the active metal surface at the pit bottom and prevents the repassivation process [108]. Furthermore, a critical potential for salt-film formation Esf can be determined, below which no salt film can form. The critical potential for salt-film formation amounts to 0.19 Vsce and 0.08 Vsce for 302 stainless steel and nickel, respectively, and does not depend significantly on the bulk concentration. For 304 SS, the lowest potential, Epi, where pit initiation occurs and current transients (metastable pitting) have been found by microelectrochemical measurements [11] coincides with the critical potential for salt-film formation, Esf, as well as with the photoelectrochemically determined flatband potential Efb [129], below which no chloride adsorption takes place. Further investigations are necessary to clarify whether these values coincide accidentally or a possible correlation exists. The above discussion of pitting corrosion leads to the following concept of localized corrosion. At potentials even below the pitting potential but above Epi a nonzero

REFERENCES

probability of pit initiation exists, depending largely on the type, size, amount, and chemical properties of these local heterogeneities. Once a pit begins to grow in the presence of aggressive anions, such as chlorides a thin salt film may be formed after a transition time. As long as the salt film is present, stable dissolution occurs, and the growth rate is diffusion controlled. If the potential is then lowered, pits may continue to grow if the composition of the pit electrolyte, the current density, and the potential at the pit bottom are such that repassivation is not possible. The growth process in this case is then mixed ohmic/charge transfer controlled [98, 105, 107]. Under these conditions, the ohmic potential drop inside and outside of the pit may play an important role. G. SUMMARY Localized corrosion is an important but complex problem responsible for many corrosion failures in engineering applications. The local breakdown of passivity of commercially available engineering materials, such as stainless steels, nickel, or aluminum, occurs preferentially at sites of local heterogeneities, such as inclusions, second-phase precipitates, or even dislocations. The size, shape, distribution, as well as the chemical or electrochemical dissolution behavior (active or inactive) of these heterogeneities in a given environment, determine to a large extent whether pit initiation is followed either by repassivation (metastable pitting) or stable pit growth, Microelectrochemical techniques, combined with statistical evaluation of the experimental results allow to gain more insight into the mechanisms of these processes. In addition to the local activation or pit initiation process, the stability of the passive film is decisive for the corrosion resistance of passive metals and alloys. Fast and effective repassivation, necessary for highly corrosion-resistant alloys, may only occur if highly stable films are formed during repassivation. Therefore, further investigations should be focused not only on the initiation of localized corrosion, but also on the stability of passive films. The stability of passive films is often reflected by the semiconductive properties of these films. Therefore electrochemical impedance spectroscopy, photoelectrochemical methods as well as in situ analytical techniques are very valuable tools to study the chemical and electrochemical behavior of these passivating oxide films. REFERENCES 1. Z. Szlarska-Smialowska, Pitting Corrosion of Metals, NACE, Houston, TX, 1986. 2. A. J. Sedriks, Corrosion of Stainless Steels, Corrosion Monograph Series, Wiley, New York, 1996. 3. H. Horvath and H. H. Uhlig, J. Electrochem. Soc., 115, 791 (1968).

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63. K. Osozawa and N. Okato, in First Soviet-Japanese Seminar on Corrosion and Protection of Metals, Moscow, Y. Kolotyrkin (Ed.), NAUKA, 1979, p. 229. 64. K. Sugimoto, M. Seto, S. Tanaka, and S. Hara, J. Electrochem. Soc., 140, 1586 (1993). 65. S. Tanaka, N. Hara, and K. Sugimoto, Mater. Sci. Eng., A198, 63 (1995). 66. P. Schmuki, M. B€ uchler, S. Virtanen, H. B€ ohni, H. M€ uller, and L. J. Gauckler, J. Electrochem. Soc., 142, 3336 (1995). 67. S. Virtanen, P, Schmuki, H. B€ ohni, P. Vuoristo, and T. M€antyl€a, J. Electrochem. Soc., 142, 3067 (1995). 68. S. Virtanen, P. Schmuki, A. J. Davenport, and C. M Vitus, J. Electrochem. Soc., 144, 198 (1997). 69. S. Virtanen, P. Schmuki, M. B€ uchler, and H. B€ ohni, Analusis, 25, M 22 (1997). 70. S. Virtanen, P. Schmuki, H. S. Isaacs, M. P. Ryan, L. Oblonsky, and H. B€ ohni, in Passivity and Its Breakdown, P. M. Natishan, H. S. Isaacs, M. Janik-Czachor, V. A. Macagno, P. Marcus, and M. Seo (Eds.), Electrochemical Society, Pennington, NJ, 1998, p. 171. 71. E. Otero, R. Lizarbe, and S. Feliu, Br. Corros. J., 13, 82 (1978). 72. W. H. Ailor, Br. Corros, J., 1, 237 (1966). 73. H. Godard, W. P. Jepson, M. R. Bothwell, and R. L. Kane, The Corrosion of Light Metals, Wiley, New York, 1967. 74. T. E. Wright, H. P. Godard, and I. H. Jenks, Corrosion, 13, 481t (1957). 75. H. S. Campbell, J. Inst. Met., 93, 97 (1964). 76. D. O. Sprowls and M. E. Carlisle, Corrosion, 17, 125t (1966). 77. P. M. Aziz and H. P. Godard, Ind. Eng. Chem., 44, 1791 (1951). 78. F. C. Porter and S. E. Hadden, J. Appl. Chem., 3, 385 (1953). 79. W. A. Bell and H. S. Campbell, Br. Corros. J., 1, 72 (1966). 80. G. Butler, P. Stretton, and J. G. Beyon, Br. Corros. J., 7, 168 (1972). 81. N. D. Tomashov, G. P. Chernova, and O. N. Markova, Prot. Met., 7, 85 (1971). 82. I. L. Rosenfeld and I. S. Danilov, Corros. Sci., 7, 129 (1967). 83. F. Hunkeler, On the Pitting Mechanism of Aluminum with Special Emphasis on Pit Growth Kinetics (in German), Ph.D. Thesis, ETH Z€ urich Nr. 6663, 1980. 84. H. Kaesche, Z. Phys. Chem., NF 34, 87 (1962). 85. H.-J. Engell and N. D. Stolica, Arch. Eisenh€ uttenw., 30, 239 (1959). 86. J. Tousek, Corros. Sci., 15, 147 (1975). 87. J. Tousek, Werkst. Korros., 25, 496 (1974). 88. G. Herbsleb and H.-J. Engell, Z. Phys. Chem., 215, 167 (1960). 89. G. Herbsleb and H.-J. Engell, Werkst. Korros., 17, 365 (1966). 90. J. Wenners, Dissolution Mechanisms within Pits on Iron and Nickel (in German), Ph.D. Thesis, Freie Universit€at Berlin, 1977. 91. F. Hunkeler and H. B€ ohni, Werkst. Korros., 32, 129 (1981). 92. F. Hunkeler and H. B€ ohni, Corrosion, 37, 645 (1981).

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14 STRESS CORROSION CRACKING R. N. PARKINS* University of Newcastle upon Tyne, Newcastle upon Tyne, UK

A. B. C. D.

Introduction Environmental conditions for SCC Role of stress in SCC Prevention and control of SCC D1. Metallurgical approaches D2. Environmental approaches D3. Stress control References

A. INTRODUCTION Metals and alloys subjected to tensile stresses and exposed to certain environmental conditions may develop cracks that would not occur in the absence of either of those controlling parameters. Not all environments that are corrosive to a particular metal promote stress corrosion cracking (SCC), but even some apparently innocuous substances, such as water, may induce cracking in some materials [1], the composition and structure of which can play a critical role in the incidence or otherwise of SCC. The manifestation of the cracks may create the impression of brittleness in the metal because the cracks often propagate with little attendant deformation, although almost invariably the properties of the metal conform to ductility specifications. The incidence of SCC appears to have increased over the last few decades, possibly because as the problem of general corrosion has been overcome, by control of environmental factors and the use of inherently more corrosion-resistant

Deceased.

materials, the probability of more localized forms of corrosion has increased. Those trends have been accompanied by moves toward higher operating stresses, deriving from the more efficient use of materials, and the more extensive use of welding as a method of fabrication. The result gives rise to residual stresses, as do other methods of fabrication involving inhomogeneous deformation, and unless such stresses are relieved, they can promote SCC, the incidence of which is probably greater from the presence of residual than from operating stresses alone. An example of SCC resulting from residual welding stresses is shown in Figure 14.1. Here liquor seeping from the cracks reveals the positions of the latter, which are associated with a longitudinal weld where the principal tensile residual stresses would be parallel to the weld and of the order of the yield stress in this nonstress relieved structure [2]. The low rate of growth of those cracks, the facts that they are visible, and because the operating pressure is relatively low, would allow appropriate action to be taken. But invariably that is not the case since the cracks may grow undetected. This was the case in the failure shown in Figure 14.2, which is from a buried high-pressure gas pipeline that failed in service and where the main fracture is apparent from the separation of the fracture surfaces toward the top of the photograph. But secondary cracks, emanating from the soil side, are also apparent [3]. These cracks are not associated with a weld and the operating stresses undoubtedly played a major role in the failure. Both Figures 14.1 and 14.2 show the presence of multiple cracks, which would be seen to be even more prevalent if examined at higher magnifications. Multiple cracking frequently also accompanies SCC failures. Stress corrosion cracking may be associated with intergranular or transgranular paths through the metal and, in some cases, with a mixture of those modes. Examination of metallographic sections from the steels involved with Figures 14.1

Uhlig’s Corrosion Handbook, Third Edition, Edited by R. Winston Revie Copyright 2011 John Wiley & Sons, Inc.

171

172

STRESS CORROSION CRACKING

FIGURE 14.1. Stress corrosion cracks in shell of coal gas liquor recirculating tank [2].

FIGURE 14.2. Secondary SCC on outer surface of high-pressure gas transmission pipeline near origin of SCC failure [3].

and 14.2 revealed intergranular cracks, with a typical example in Figure 14.3. An example of an essentially transgranular crack, in a different pipeline steel from that involved with Figure 14.3 and exposed to a different environment [4], is shown in Figure 14.4. Figure 14.5 shows an example of

mixed-mode cracking in an arsenical a brass exposed to water containing a small amount of sulfur dioxide [5]. In the same system, cracks may initiate in the intergranular mode, then change to transgranular as the stress concentration increases with increasing crack depth, and even revert to intergranular if the crack bifurcates and reduces the stress concentration. Since the most common mode of failure of metals is transgranular (e.g., by overload or fatigue), but not high-temperature creep, intergranular SCC is sometimes regarded as unusual and has often been associated with the collection of segregates or precipitates at grain boundaries. That is undoubtedly so in some cases, of which the propensity for intergranular SCC in sensitized austenitic stainless steels, due to the precipitation of chromium carbides in boundaries and the impoverishment of the surrounding material in chromium is probably the best-known instance [6]. However, as the examples quoted above involving pipeline steels indicate, a change in environment may result in a change of crack path, while the examples involving brass indicate that a change in the stressing conditions may also result in a change in cracking mode. Thus, while electrochemical heterogeneity due to segregates or precipitates at grain boundaries may play a role in some instances of intergranular SCC, the incidence of the latter or otherwise is dependent on the environmental and stressing conditions. The interactions between environmental and stressing conditions upon the paths of cracks indicate the possible problems in deducing some mechanistic hypothesis for SCC based upon limited studies of a particular system. Mechanisms of SCC have been, and continue to be, widely discussed. Most suggested mechanisms invoke either a process of localized embrittlement of the metal in the vicinity of the crack tip or of localized dissolution in that region [7]. There are variations on both of these themes, so that, for example, while hydrogen ingress may result in the embrittlement of some metals, and the mechanism of such embrittlement is a matter of debate, the formation of films with certain properties may induce cleavage which, once initiated, continues to advance into the underlying metal before arresting. Such film-induced cleavage may involve dealloyed layers in some alloys, but those layers may play the same role as oxide films, in preventing lateral dissolution on the crack sides and concentrating dissolution in the crack tip region, when bare metal is exposed by straining in the adjacent metal. Such strains may be associated with transgranular cracking where bare metal slip steps emerge through films in the crack tip region, or they may induce intergranular cracking, especially if electrochemical heterogeneity causes preferential dissolution at the grain boundaries. It is possible that with some environments, both dissolution and hydrogen ingress may facilitate crack growth. It is not surprising that mechanistic aspects of SCC are the subject of almost annual conferences, but for our present purposes the emphasis is on assisting the practicing engineer predict those combinations of metal and environment that promote SCC and how they may be

INTRODUCTION

FIGURE 14.3. Intergranular cracks in pipeline steel [3].

FIGURE 14.4. Transgranular cracking in different pipeline steel from that of Figure 14.3 produced by different environment [4].

FIGURE 14.5. Mixed-mode SCC in a brass exposed to H2O þ SO2 [5].

173

174


controlled or avoided in particular engineering structures. Complete mechanistic understanding of SCC would assist in achieving such goals, but in the absence of such understanding it is still possible to make some advances in relation to practical control.

B. ENVIRONMENTAL CONDITIONS FOR SCC Until about four decades ago, it was thought that SCC occurred only in particular alloys exposed to a few very specific environments (e.g., ferritic steels exposed to hydroxides or nitrates, brasses exposed to ammoniacal environments, and austenitic stainless steels or aluminum alloys to chlorides). While the concept of solution specificity remains, since not all environments corrosive toward a particular alloy promote SCC, the number of environments that will promote that mode of failure has increased considerably in recent times. It is impossible to list here all of the environments that have been shown to promote SCC in the commonly used alloys, but, in terms of predicting whether or not a particular metal–environment combination is likely to suffer SCC, certain circumstances need to obtain. Where crack growth is by a localized dissolution process, potent solutions will need to promote a critical balance between activity and passivity since a highly active condition will result in general corrosion, while a completely passive condition cannot lead

to SCC. For most engineering alloys, inactivity at exposed surfaces is due to the presence of adherent oxide films on those surfaces. Then, it is not surprising that the alloys of high inherent corrosion resistance (e.g., austenitic stainless steels that readily develop protective films) require an aggressive ion, such as a halide, to promote SCC. Alternatively, metals of low inherent corrosion resistance (e.g., C steels) require the presence of an environment that is itself partially passivating for SCC. Such steels can fail in solutions of anodic inhibitors (e.g., nitrates, hydroxides, carbonates, or phosphates). The SCC is influenced not only by the presence of particular ions and their concentration but also by the electrode potential, so that cracking occurs only within certain potential ranges for particular metal–environment combinations. Figure 14.6 shows the cracking domains for ferritic steels exposed to various environments (also involving different temperatures, another parameter that can influence SCC), and pH and potential dependent cracking domains have been shown to exist for other alloys. The boundaries of such domains would be expected to relate to particular reactions involving film formation. For each of the systems shown in Figure 14.6, the upper boundaries correspond to the formation of Fe2O3, with ductile failure in slow strain rate tests at potentials above the cracking domain. There are two exceptions in that at potentials high enough to form Fe2O3, nitrates and high-temperature water promote cracking from within pits. The initiation of cracks from pits has been

FIGURE 14.6. SCC potential–pH domains for ferritic steels in various environments, the latter at a variety of temperatures, together with equilibrium potentials for discharge of hydrogen at 25 C.

ENVIRONMENTAL CONDITIONS FOR SCC

observed in many systems and while that has sometimes been related to stress intensification at the base of pits, it is probably at least as likely to relate to changes in the composition of the solution and potential within the pits. With some bulk solutions, lowering of the pH and potential will occur in pits, so that the SCC domain for the bulk solution may be irrelevant and indeed the mechanism of cracking may change. Thus, Figure 14.6 shows that the cracking domains are above the equilibrium potentials for hydrogen discharge from the solutions of various pH values, but even the buffered solutions (OH, CO3/HCO3, C2H3O2, and PO4), which do not cause pitting at the concentrations involved with Figure 14.6, can promote hydrogen-related cracking if the potential is reduced sufficiently for hydrogen discharge. However, where acidification of the localized environment within a pit occurs, hydrogen discharge is likely to be facilitated. While it is possible to predict cracking domains from thermodynamic data for some of the systems to which Figure 14.6 refers, there are difficulties with such predictions when the bulk environment is changed within a pit or crack enclave. Moreover, even successful predictions have followed from experimental determination of the cracking domain boundaries and, at present, there is no theoretical approach to determining whether or not a particular metal– environment combination, for which there is no prior experience, will promote SCC. However, there are some relatively rapid experimental approaches to determining the potency of systems for SCC, where the bulk environment is itself the potent solution. The latter may derive from the bulk solution, which itself may be incapable of promoting SCC, by concentration in a crevice or at a heat transfer surface. In these circumstances predictability may be more difficult, unless the possibility of such concentration is recognized. The instances of SCC in pipelines mentioned in the context of Figures 14.2–14.4 illustrate the point. The transgranular cracking (Fig 14.4) is due to the presence of a dilute groundwater containing CO2 (pH 6.5), while the intergranular cracking (Figs. 14.2 and 14.3) is due to the generation of a relatively concentrated carbonate–bicarbonate solution derived from groundwater but concentrated due to the flow of cathodic current at the pipe surface and ion transport, as well as heat transfer in the crevice between the pipe and a disbonded coating. Where crack advance is by dissolution, the crack tip must be active, but as the crack advances the crack sides must become relatively inactive, otherwise the sides will extend laterally and the geometry will be changed to that of a pit. Figure 14.4 shows evidence of corrosion on the crack sides increasing in moving away from the tip, so that the amount of activity acceptable on the crack sides will depend on the rate of crack growth. Transitions from electrochemically active to relatively inactive behavior may be expected to be reflected in the current response of the bare metal exposed to the appropriate environment. Thus, dissolution will be associated with the passage of relatively high anodic current densities, but

175

with the passage of time this current will decay if filming occurs. Very rapid rates of decay are not likely to permit much dissolution and so are not likely to be indicative of conditions conducive to SCC. Very slow rates of decay are more likely to be indicative of insufficient development of inactivity to retain crack geometry. There are various techniques for measuring these features, including scratching or rapidly straining electrodes previously filmed at particular potentials. A convenient way of anticipating the range of potentials in which SCC may occur is through potentiodynamic polarization curves determined at different sweep rates [8]. If the potential of an initially film-free surface is rapidly changed ( 1 V/min) over an appropriate range, then the currents passed at the surface will indicate ranges of potential in which relatively high anodic activity is likely. The rapid sweep of the potential range has the object of minimizing film formation. If the experiment is repeated at a slow rate ( 10 mV/ min) of potential change allowing filming, comparison of the two curves will indicate ranges of potential within which high anodic activity in the film-free condition reduces to insignificant activity at the slow sweep rate, thereby identifying the range of potentials in which SCC is likely. The method correctly anticipates SCC of ferritic steels in a number of very different environments, but is only applicable where airformed oxide films can be reduced. Thus for metals with very stable films one resorts to scraping or rapid straining as a means of creating bare metal. For systems where cracks may be initiated from pits, measurement of the pitting potential can give an indication of the minimum potential for cracking, while for systems where hydrogen-induced cracking is possible calculation of the equilibrium potential for hydrogen discharge, from the pH of the solution, will give an indication of the highest potential for such cracking. However, in the application of any of these approaches to estimating potentials for cracking it is well to remember that the surfaces in plant may differ appreciably in terms of potentials from those measured in the laboratory, the latter usually on carefully polished surfaces, in the same environment. Consequently, a laboratory SCC test at open circuit potential may not give an adequate indication of the propensity for cracking in a particular system. A few additional tests at controlled potentials, defined by some appropriate electrochemical measurements, are therefore necessary. This is especially so if cracking does not occur at open circuit. In relation to the stressing of specimens for assessing the propensity for SCC in any system, slow strain rate tests (SSRTs) offer a rapid method of arriving at a result [9, 10]. These, which may be conducted on initially plain or precracked specimens, are simply tensile tests conducted at relatively slow strain rates, typically of the order of 10 6/s for steels and Cu or Ni alloys, and 10 5/s for Ti or Mg alloys. Failure will usually occur in 1 or 2 days and various parameters may be employed for quantifying the results, particularly those related to ductility, although

176


fractographic or metallographic examination of the failed samples should be conducted. Obviously, the stressing and/or straining conditions in SSRTs are beyond those likely to be experienced in service. Strain rate effects in the latter are considered later, but where the objective of the laboratory test is to give a “go–no go” result quickly, then SSRTs are particularly useful.

C. ROLE OF STRESS IN SCC It is usual to consider the role of stress, or the stress intensity factor in the case of precracked specimens, in terms of its influenceuponthetimetofailure inagivensystem.Figure14.7 shows some typical results [11], which indicates that there is a stress, often referred to as the threshold stress, below which failure does not occur in an extended test time. The threshold stress is not only a function of alloy composition and structure, but also of the environmental conditions, including solution composition, potential, and temperature, so that it is not a unique property of a material in the sense of a yield or tensile strength. Moreover, while the threshold stress is sometimes defined as the stress below which cracking does not occur, this is not necessarily so and it is better defined as the stress above which total failure occurs, since for some systems cracks have been shown to initiate below die threshold but to cease to propagate after some growth. It is difficult to explain why cracks should cease to propagate on any stress-based argument, since stress concentration or intensification would be expected to increase with crack growth under constant load. A

feasible explanation is that it is not stress per se but the strain rate it engenders that is the controlling factor, and that cracks cease to propagate when the crack tip strain rate falls below some critical value related to the rate of film growth. Such an explanation is consistent with the influence of the relative times at which the stress and the environmental conditions for cracking are established, creep at constant load prior to the establishment of the environmental conditions delaying or preventing cracking in laboratory tests. Obviously, it is also consistent with the demonstration of the influence of applied strain rate upon cracking for a wide variety of combinations of metal and environment, for some of which it has been shown that sufficiently slow rates of straining can result in ductile failure without impairment of the tensile strength of the alloy despite its exposure to a potent cracking environment. There is another important consequence of the significance of strain rate in facilitating crack growth and it relates to cyclic, as opposed to static, loading. Hysteresis effects are well known to accompany cyclic loading and the cyclic stress–strain curves of materials often fall appreciably below their monotonic loading counterparts, reflecting the fact that cyclic loading facilitates microplastic deformation [12]. It follows from such observations that load cycling may produce SCC at significantly lower stresses than those needed with static loading. Figure 14.8 shows this to be so for a pipeline steel exposed to a carbonate–bicarbonate solution by comparison of the data for static and cyclic loading conditions [13]. In the absence of cyclic loading (DK ¼ 0), the threshold stress intensity factor, KIscc , is 21 MN/m3/2, and cyclic loading at the frequency of 11 Hz did not alter that

FIGURE 14.7. Initial stress–time to failure curves for ferritic steels with different Ni contents in boiling 4N NH4NO3. Data points for only one steel are shown but those give an indication of the scatter for each steel [11].

PREVENTION AND CONTROL OF SCC

177

FIGURE 14.8. Modified Goodman diagram indicating loading parameters for which intergranular or transgranular cracking were observed in pipeline steel exposed to carbonate–bicarbonate solution at 75 C and 0.65 V(SCE) [13].

threshold, although intergranular cracking extended to DK values of 5 MN/m3/2 before a transition to transgranular cracking. With reduction of the frequency of load cycling to 0.19 Hz, not only is the purely intergranular cracking observed at higher DK values, but the threshold mean stress intensity factor is reduced to 10 MN/m3/2 at small values of DK. Similar large reductions in threshold stresses have been observed with initially plain specimens when subjected to cyclic loading at low frequencies and stress amplitudes of 20% of the mean stress. Of course, it may be argued that with cyclic loading what is being studied is corrosion fatigue, but fractographically the cyclically loaded specimens are indistinguishable from statically loaded samples, providing that the DK values are maintained below those levels that promote transgranular cracking. Indeed it may be argued that the latter mode of cracking is more typical of what may be expected in corrosion fatigue in such a system. The distinction here then is not between static and cyclic loading, or SCC and corrosion fatigue, but rather between relatively small stress amplitudes at low frequencies, sometimes referred to as ripple loading, and the much higher stress amplitudes usually involved with fatigue. Such distinctions are inevitably rather arbitrary, but there is another point to be remembered in this context, which relates to environmental influences. The environments that will promote corrosion fatigue in, say, ferritic steels are much more extensive than those that will promote SCC. A possible reason is that, as already mentioned, for SCC the environment needs to have characteristics that assist in the retention of crack geometry by filming of crack sides, but with large amplitude stresses mechanical crack sharpening occurs, so

that there are less stringent requirements for the properties of the environment.

D. PREVENTION AND CONTROL OF SCC Since the incidence of SCC requires a susceptible alloy to be exposed to a specific environment at stresses above some limiting value, it follows that control of the problem may be through manipulation of any or all of these three parameters. Ideally, approaches to prevention should begin with the selection of a resistant alloy, which is the most usual approach, followed by consideration of possible modification of the stress or environment, with variations within these three themes as outlined in Figure 14.9 [14]. However, it sometimes happens that SCC occurs in an existing plant when such had not been anticipated at the design stage or, indeed, that a susceptible material had to be used because of other considerations, in which cases the approaches to prevention are restricted. Whatever approach is pursued, it is likely to depend on experience or laboratory test data, and it is important to realize that alloy susceptibility is not simply a function of alloy composition or structure but also of the environmental conditions. Thus, Figure 14.7 shows the beneficial effects of sufficient Ni additions to a ferritic steel upon resistance to cracking, as measured by the threshold stress, in a nitrate environment. However, if the same steels are exposed to a boiling MgCl2 solution, then the Ni additions have the effect of increasing cracking susceptibility, although much higher ( 50%) additions may promote immunity, based upon data for austenitic steels containing some 18–20% Cr [15].

178


FIGURE 14.9. Approaches to SCC control.

D1. Metallurgical Approaches There are many published studies of the manipulation of the composition or structure of alloys to control SCC [16]. Where intergranular cracking is due to segregation at the grain boundaries, control of the segregant affords a means of controlling SCC. Thus, restricting the carbon content of stainless steels to a maximum of 0.02%, whereby weld sensitization is avoided, can markedly improve cracking resistance [17]. There are other examples of relatively small changes in alloy composition markedly influencing cracking resistance (e.g., in aluminum alloys [18]), but in many instances, relatively large additions of alloying elements are necessary to achieve significant improvements in resistance. The point may be illustrated by data for additions to ferritic steels in relation to their SCC susceptibility in NO3 or OH solutions [11]. It has already been mentioned that cracking is a function of potential for a given environment, and determination of that relationship in SSRTs affords a means of comparing the effects of alloying additions. The susceptibility in an SSRT may be expressed in terms of the ratio of the time to failure in the potent solution to that in an inert environment (oil) at the same temperature, a ratio of 1 indicating no susceptibility and increasing departure from 1 indicating increasing susceptibility. A plot of that ratio against potential bounds an area within which SCC occurs, that area being a measure of the stress corrosion index (SCI), and reductions or increases in that area indicate beneficial or deleterious effects, respectively, for alloying additions.

Multiple regression analysis of the data resulted in the following equations for the relative effects of various alloying additions upon SCI in NO3 and OH solutions: SCIðNO3 Þ ¼ 1777 996% C 390% Ti 343% Alð 132%MnÞ 111% Cr 90% Mo 62% Ni þ 292%Si SCIðOHÞ ¼ 105 45% C 40% Mn 13:7% Ni 12:3% Cr 11% Ti þ 2:5% Al þ 87% Si þ 413% Mo The first constant on the right-hand side of those expressions reflects the greater propensity for cracking and the wider potential range involved with NO3 than OH solutions, while negative coefficients for the various alloying elements indicate beneficial effects and positive coefficients deleterious influences. (Where a coefficient is bracketed, the t ratio, the coefficient/standard error of the coefficient, was KISCC . In this model, the cyclic character of loading is not important. The combination of true corrosion fatigue and stress corrosion fatigue results in type C, the most general form of corrosion fatigue crack propagation behavior, Figure 16.4(e), which involves cyclic timedependent acceleration in da/dN below KISCC , combined with time-dependent cracking (SCC) above the threshold. Figure 16.5 shows examples of corrosion fatigue crack propagation behavior. Figure 16.5(a) illustrates the behavior of maraging steel exposed to 3% NaCl [18], representing Type A growth, and Figure 16.5(b) shows the behavior of high-strength type 4340 [Unified Numbering System (UNS)

CORROSION FATIGUE CRACK PROPAGATION

201

FIGURE 16.6. Corrosion fatigue of Al–Zn–Mg alloy, 7017 in natural seawater. (a) Crack growth rate as a function of DK for a range of cyclic loading frequencies, (b) The dependence of corrosion fatigue fracture morphologies in terms of cyclic loading frequency and DK. (c) Fracture morphologies in terms of crack growth rate and cyclic loading frequency. [Reprinted from N. J. H. Holroyd and D. Hardie, “Factors Controlling Crack Velocity in 7000 Series Aluminum Alloys during Fatigue in an Aggressive Environment,” Corrosion Science, 23, pp. 529, 533, 535 (1983), with permission from Elsevier Science.]

202

CORROSION FATIGUE

G43400] steel in water vapor and argon [19], representing type B growth. In Figure 16.5(a), there is a substantial corrosion fatigue effect below the static load threshold, but only for those load waveforms that include a slow deformation rate to maximum stress intensity. The solid line in Figure 16.5(b) demonstrates that time-dependent corrosion fatigue crack growth rates are accurately predicted by linear superposition of stress corrosion crack growth rates (da/dt) integrated over the load–time function for fatigue. A number of interactive variables influence the relationship between corrosion fatigue crack growth rate and stress intensity. Growth rates are affected by environmental chemical variables (e.g., temperature; gas pressure and impurity content; electrolyte pH, potential, conductivity, and halogen or sulfide ion content); by mechanical variables, such as DK, mean stress, frequency, waveform, and overload; and by metallurgical variables, including impurity composition, microstructure, and cyclic deformation mode. Time, or loading frequency, is also critical. Figure 16.6 shows the effect of loading frequency and stress intensity range on the corrosion fatigue crack growth rate and the cracking morphology of Al–Zn–Mg alloy, 7017, in natural seawater [20]. Although da/dN DK shows type B growth behavior, a simple superposition model is inappropriate for describing the corrosion fatigue crack growth rate and the effect of load frequency. The cracking morphology, intergranular or transgranular, is influenced by both the load frequency and the stress intensity factor range, and intergranular cracking can occur at a very high load frequency (70 Hz) as long as the DK values are sufficiently low. The frequency dependence of the crack velocities associated with the transition from intergranular to transgranular cracking shows a linear relationship with the square root of the loading cycle period, implying that the rate-controlling step is consistent with grain boundary diffusion of hydrogen during the loading cycle. Corrosion fatigue can be prevented by using highperformance alloys resistant to corrosion fatigue; but for most engineering applications this approach may not be practical because of the availability and cost of these alloys. In general, methods that reduce corrosion rate and/or cyclic damage can be beneficial for eliminating corrosion fatigue damage. While effective coatings and inhibitors can delay the initiation of corrosion fatigue cracks, improving surface conditions is also very useful. Compared with reducing the maximum stress level, it is often more beneficial and more cost effective to reduce the magnitude of the stress fluctuation.

REFERENCES 1. O. Devereux, A. J. McEvily, and R. W. Staehle (Eds.), Corrosion Fatigue: Chemistry, Mechanics and Microstructure, NACE-2, NACE, Houston, TX, 1972.

2. T. W. Crocker and B. N. Leis (Eds.), Corrosion Fatigue: Mechanics, Metallurgy, Electrochemistry and Engineering, ASTM Special Technical Publication 801, American Society for Testing and Materials, Philadelphia, PA, 1984. 3. R. P. Gangloff and M. B. Ives (Eds.), Environment-Induced Cracking of Metals, NACE-10, NACE, Houston, TX, 1990. 4. P. M. Scott, “Chemical Effects in Corrosion Fatigue,” in Corrosion Fatigue: Mechanics, Metallurgy, Electrochemistry and Engineering, T. W. Crooker and B. N. Leis (Eds.), ASTM Special Technical Publication 801, American Society for Testing and Materials, Philadelphia, PA, 1984, p. 319. 5. R. P. Gangloff, “Corrosion Fatigue Crack Propagation in Metals,” in Environment-Induced Cracking of Metals, R. P. Gangloff and M B. Ives (Eds.), NACE, Houston, TX, 1990, p. 45. 6. D. J. Duquette, “Corrosion Fatigue Crack Initiation Processes: A State-of-the Art Review,” in Environment-Induced Cracking of Metals, R. P. Gangloff and M. B. Ives (Eds.), NACE-10, NACE, Houston, TX, 1990, p. 45. 7. S. Suresh, Fatigue of Materials, Cambridge Solid State Science Series, Cambridge University Press, Cambridge, UK, 1991. 8. Y.-Z. Wang, R. Akid, and K. J. Miller, Fatigue Fract. Eng. Mater. Struct., 18, 293 (1995). 9. S. Suresh, Fatigue of Materials, Cambridge Solid State Science Series, Cambridge University Press, Cambridge, UK, 1991, pp. 363–368. 10. R. Kiessling, Non-Metallic Inclusions in Steel, Metals Society, London, UK, 1978. 11. Y.-Z. Wang and U. Akid, Corrosion, 52, 92 (1996). 12. D. C. Jones, “Localized Corrosion,” in Corrosion Processes, R. N. Parkins (Ed.), Applied Science Publishers, London, UK, 1982, p. 161. 13. G. P. Ray, R. A. Jaman, and J. G. N. Thomas, Corros. Sci., 25, 171 (1985). 14. D. J. Duquette and H. H. Uhlig, Trans. Am. Soc. Metals, 62, 839 (1969). 15. H. H. Lee and H. H. Uhlig, Metall. Trans., 3, 2949 (1971). 16. Y.-Z. Wing, R. W. Revie, and R. N. Parkins,“Mechanistic Aspects of Stress Corrosion Crack Initiation and Early Propagation,” Paper No. 99143, CORROSION/99, NACE International, Houston, TX, 1999. 17. A. J. McEvily and R. P. Wei, “Fracture Mechanics and Corrosion Fatigue,” in Corrosion Fatigue: Chemistry, Mechanics and Microstructure, O. Devereux, A. J. McEvily, and R. W. Staehle (Eds.), NACE-2, NACE, Houston, TX, 1972, p. 381. 18. J. M. Barsom, “Effect of Cyclic Stress Form on Corrosion Fatigue Crack Propagation Below KISCC in a High-YieldStrength Steel,” in Corrosion Fatigue: Chemistry, Mechanics and Microstructure, O. Devereux, A. J. McEvily, and R. W. Staehle (Eds.), NACE-2, NACE, Houston. TX, 1972, p. 426. 19. R. P. Wei and G. W. Simmons, “Environment Enhanced Fatigue Crack Growth in High-Strength Steels,” in Stress Corrosion Cracking and Hydrogen Embrittlement of Iron Base Alloys, R. W. Staehle, J. Hochmann, R. D. McCright, and J. E. Slater (Eds.), NACE-5, NACE, Houston, TX, 1973, p. 751. 20. N. J. H. Holroyd and D. Hardie, Corros. Sci., 23, 527 (1983).

17 FLOW EFFECTS ON CORROSION K. D. EFIRD Efird Corrosion International, Inc., The Woodlands, Texas

A. Introduction B. Flow basics related to corrosion B1. Turbulent boundary layer B2. Boundary layer disruption B3. Flow parameters B3.1. Wall shear stress B3.2. Mass transfer coefficient B3.3. Interrelationship of mass transfer coefficient and wall shear stress C. Flow regime and flow corrosion C1. Single-phase flow C2. Multiphase flow and flow regime C3. Effect of disturbed flow C4. Low-turbulence-flow corrosion Nomenclature References

A. INTRODUCTION The effect of flow on corrosion is complex and varied and is dependent on both the chemistry and physics of a system. The effect of chemistry is typical of most corrosion mechanisms except that flow has a significant effect on the kinetics of the corrosion reactions and the phases participating in the corrosion reaction. These effects are related to the influence of flow on the movement, distribution, and mixing of fluids in the flowing system. The key variable defining the effect of flow on corrosion is turbulence. High turbulence can result in flow-induced corrosion, erosion–corrosion, or cavitation. Low turbulence can result in corrosion in a separated water phase and allows the

occurrence of corrosion under deposits and/or in separated liquid water. The emphasis is on flow in pipes since pipe flow is the location of the vast majority of flow-induced corrosion problems. “Flow-induced corrosion” is the term used to describe the increase in corrosion resulting from high fluid turbulence due to the flow of a fluid over a surface in a flowing single or multiphase system. “Underdeposit corrosion” is the term used to describe the increased corrosion occurring in a separated water phase beneath deposits of nonmetallic solids on a metal surface resulting from low-flow turbulence. A clear understanding of the difference between flowinduced corrosion and “erosion–corrosion” aids in the discussion. The two are not the same, as indicated by the definitions of these terms in the Glossary. Flow-induced corrosion is the increase in corrosion resulting from increased fluid turbulence intensity and mass transfer as a result of the flow of a fluid over a surface. Erosion–corrosion is the increased corrosion due to the physical impact on a surface causing mechanical damage. The impact can be from solid particles entrained in a liquid or gas phase or liquid droplets entrained in a gas phase. The effect of erosion on corrosion is covered in Chapter 18, Erosion–Corrosion in Single- and Multiphase Flow.

B. FLOW BASICS RELATED TO CORROSION Corrosion is a surface phenomenon, and, as such, what goes on at the metal surface has a profound effect on corrosion. Many aspects of fluid dynamics relate to and define the interactions of a fluid with the surface that are important to corrosion.


203

204

FLOW EFFECTS ON CORROSION

When a fluid flows over a solid surface, the flow is characterized as either laminar or turbulent. In most situations where the effect of fluid flow on corrosion is important, the flow is turbulent; so the physical structure of turbulent flow is a primary consideration. A number of possibly unfamiliar terms introduced in this chapter are defined in the Glossary. The definitions of symbols used in equations are given in the Nomenclature at the end of this chapter. B1. Turbulent Boundary Layer Fully developed turbulent flow consists of a turbulent core where the mean velocity is essentially constant and a boundary layer of varying fluid velocity near the solid–fluid interface. The majority of the changes in fluid stress characteristics, turbulence, mass transfer, and fluid interaction with the wall occur within this boundary layer. All mean and turbulent velocity components of the flow must go to zero at the wall and must be very small in the immediate vicinity of the wall. Therefore, the components of all turbulent stresses must also go to zero at the wall, leaving only the viscous stresses of laminar flow to act on the wall. This layer near the wall within the boundary layer, where viscous forces dominate over turbulent forces and flow is laminar-like, is termed the viscous sublayer. There is a transfer of turbulent energy from the outer layer to the viscous sublayer and a turbulent diffusion of energy from the laminar sublayer to the outer layer at approximately equal rates for established flow. A diagram of these layers and the transfer processes is shown in Figure 17.1 [1]. The calculated boundary layer thickness generally referred to in the corrosion literature is related to and representative of the physical boundary layer but is actually a defined thickness that is a linear approximation [2]. This approximation for the diffusion boundary layer (dd) and the hydrodynamic boundary layer (dn) is shown graphically in Figure 17.2 [3]. The approximation assumes that the flowing fluid can be broken down into two parts: (1) a thin boundary layer near the solid surface and (2) the bulk solution, and that

the transition occurs abruptly, defining the boundary layer thickness. The basic assumption is that momentum and mass transport occur by viscous flow and diffusion within the boundary layer, as measured by the mass transfer coefficient (k) and the shear stress (t), and by turbulent flow and convection outside the boundary layer [2]. The turbulent boundary layer is further resolved into three regions: an outer “logarithmic” region (30 < y þ < 100), an intermediate “buffer” region (5 < y þ < 30), and, very close to the solid wall, a “viscous” region (y þ < 5), where y þ is the dimensionless viscous length perpendicular to the wall, defined as y þ ¼ yut =n [4]. A diagram of these layers along with an idealized plot of the normalized turbulent velocity variation is shown in Figure 17.3 [4]. Experimental details of the variation in boundary layer turbulence with distance from the wall are given by Kline et al. in their extensive study into the structure of the turbulent boundary layer [3]. The nature of the turbulent boundary layer dictates that this is where the processes that control corrosion and film formation will occur. The majority of the movement of corroding species to, and of corrosion products from, the wall and chemical reactions with the wall must occur in this region. Therefore, any disturbance in the turbulent boundary layer must be a primary factor affecting the corrosion process. B2. Boundary Layer Disruption Disruption of the boundary layer in turbulent flow occurs primarily by the formation of turbulent bursts and sweeps. The turbulent burst is an ejection of fluid from the wall, which also causes fluid to impinge on the wall by the simultaneous formation of sweeps, or movement of fluid toward the wall. Turbulent bursts and sweeps occur through the formation of vortices and the lift-up of wall streaks [5]. One process of vortex formation and evolution into a turbulent burst proposed by Praturi and Brodkey [6] is illustrated in Figure 17.4. The vortex moves with the flow and is increasingly tilted by the mean shear in the viscous

Outer Edge of Hydrodynamic Boundary Layer Turbulent Stress Dominant

Outer Layer

Energy transfer inward by Reynolds Stress

Viscous Sublayer

Diffusion of turbulent energy from sublayer

Conversion into turbulent energy and dissipation

WALL FIGURE 17.1. Basic structure of turbulent boundary layer and associated mechanical energy flow within the boundary layer [1].

FLOW BASICS RELATED TO CORROSION

205

1E-1

1.2

δd δν

3E-2

V1 0 Concentration, mol/l

5E-2

2E-2 1E-2 5E-3

0.8 0.6

0.4 C B,W 0.2

δν d

3E-3

C B,0

0

0

0.2

δd

2E-3 1E-3 2E+4

3E+4

5E+4 Re

7E+4

0.4

δν

0.6 0.8 Distance, y

y

1

1.2

1E+5

FIGURE 17.2. Relative boundary layer thickness (d) for concentration (c) and velocity (n) profiles in a turbulent flowing fluid (solid line ¼ true profile; dotted line ¼ linear approximation) [3].

100

vx+ = 2.5 ln y+ + 5.5

Logarithmic Region Hydrodynamic Boundary Layer

y+ 30

Buffer Region 5

Viscous Region

vx+ = y+ 0

WALL

v x+

Diffusion Boundary Layer

FIGURE 17.3. Structure of the turbulent boundary layer with an idealized plot of the normalized turbulent velocity variation with distance from the wall, expressed as dimensionless viscous length [4].

region (stages 1–4). At some point, the vortex becomes unstable and a strong ejection occurs along with a rapid sweep of fluid into the viscous region (stage 5). It is evident that severe pressure fluctuations must occur with the formation of a turbulent burst and during the rapid ejection of fluid. A slightly different model of the bursting process is given by Often and Kline, where a description of the complete burst cycle is given [7]. In this model, bursting is associated with wall streaks and stretched and lifted vortices. All of the current theories of wall turbulence agree in the basic structure of the turbulence, that is, the existence and interaction of turbulent bursts, ejections, sweeps, and wall streaks. The differences lie in the mechanisms of formation of these structures and the details of their interaction.

An excellent review of the present understanding of the turbulent boundary layer structure is given by Robinson [8]. As a result of this review, Robinson concluded that coherent motions exist in the viscous region, that they consist of elongated, unsteady regions of high- and low-speed streaks, and that sweep motions dominate the viscous region. The majority of the coherent flow structures occur in the buffer region of the turbulent boundary layer, with only sweeps and ejections penetrating into the viscous layer [8]. Laser Doppler measurements have clearly demonstrated the presence of turbulence in the viscous region near the wall as shown in Figure 17.5 [4]. The relationship of turbulent structures in the boundary layer, specifically near wall bursting activity, and fluctuations

206

FLOW EFFECTS ON CORROSION Boundary Layer Thickness

Stage 1

High Speed Front

y

Transverse Vortex

x

Low Speed Region

Stage 2

Ejection Inflow (start)

Stage 3

Outflow

Ejection

Inflow

Stage 4

Circulation Outflow

New Vortex Ejection

Inflow (end)

Stage 5

Ejection

WALL

FLOW DIRECTION

FIGURE 17.4. Progression of near-wall flow illustrating the formation of shear layer vortices and evolution into turbulent bursts [6].

Re = 8200 y+ = 3.8 Intensity

Velocity

FIGURE 17.5. Laser Doppler measurements at yþ ¼ 3.8, showing turbulence inside viscous region for water flowing in rectangular pipe. The large peak is the main velocity component, and the small peaks represent instantaneous velocity fluctuations [4].

in the wall shear stress and wall pressure were measured by Thomas and Bull [9]. A diagram of their results is given in Figure 17.6. This work provides evidence of significant variations in wall shear stress related to the boundary layer turbulence generation process that could have a major effect on corrosion processes. B3. Flow Parameters Fluid flow must be expressed in terms broadly related to flow parameters that are common to all hydrodynamic systems and that effectively define the interaction of the fluid with the metal surface. The parameters that best fit this requirement are wall shear stress (tw) and mass transfer coefficient (k). These parameters are calculated from empirical equations developed to characterize fluid flow.

B3.1. Wall Shear Stress. Wall shear stress is a direct measure of the viscous energy loss within the turbulent boundary layer and is related to the intensity of turbulence in the fluid. It is defined as the isothermal pressure loss in a moving fluid within an incremental length due to fluid friction as a result of contact with a stationary wall. The mathematical definition of wall shear stress, t w, is as follows. [2] The total shear stress, t, in a fluid moving past a fixed wall is the sum of the viscous and turbulent stresses, expressed as qU ux uy t¼n qy

ð17:1Þ

where n ¼ kinematic viscosity (m2/s) U ¼ mean velocity (m/s) y ¼ direction ? to the surface (m) ux, and uy ¼ fluctuating velocity component (m/s) in the x and y direction The Reynolds stresses (uxuy) go to zero at the wall, leaving only the viscous stress in the fluid. The wall shear stress, t w, is defined as this viscous shear stress at the wall (y ¼ 0), expressed as

qU tw ¼ n qy

ð17:2Þ y¼0

B3.2. Mass Transfer Coefficient. Mass transfer defines the chemical and electrochemical effects of fluid flow. It relates to the transport of reactive chemical species from the bulk solution to and from the solid surface through the boundary layer. It defines the movement of corrosive species to the metal surface and the movement of corrosion products away

FLOW REGIME AND FLOW CORROSION

Sweep

Lift-up

207

New Streak

Break-up

τw pw FIGURE 17.6. Wall shear stress and wall pressure fluctuations in relation to location of near-wall bursting activity in turbulent boundary layer [9].

from the metal surface. The diffusion layer, in which the mass transfer occurs, is roughly defined by the diffusion mass flux density (Nb,y) as expressed in the Nernst diffusion model [10]: Nj;y

dCj Cj;0 Cj;y ¼ Dj ¼ Dj dy y dd

ð17:3Þ

where Nj,y ¼ mass flux density of species j in the y direction (mol/m2 s) Dj ¼ diffusion coefficient for species j (m2/s) Cj ¼ concentration of species j (mol/L) dd ¼ diffusion boundary layer thickness (m) The mass transfer coefficient, kd, is defined as the proportionality between the mass flux density and the concentration gradient, expressed as kd ¼

Nj;y Cj;0 Cj;y

ð17:4Þ

Substituting the definition of the flux density from Eq. (17.3), the mass transfer coefficient is defined as kd ¼

Dj dd

0:5 tw kd r

ð17:6Þ

where r is fluid density (kg/m3). This relationship is more precisely defined as 0:5 tw 2 kd ¼ 17:24 Sc =3 r

for Sc > 100

ð17:7Þ

where Sc is the Schmidt number. C. FLOW REGIME AND FLOW CORROSION C1. Single-Phase Flow Single-phase flow consists of only one phase, gas or liquid. In single-phase aqueous systems there is a flow pattern that consists of a turbulent core and a boundary layer near the solid–fluid interface as described earlier. All of the mass and momentum transfer to the boundary layer occurs through this turbulent core. The equations and relationships discussed in the sections on wall shear stress and mass transfer are directly applicable to the single-phase flow situation.

ð17:5Þ

B3.3. Interrelationship of Mass Transfer Coefficient and Wall Shear Stress. Mass transfer is related to wall shear stress in that changes in flow parameters that affect one result in changes in the other. The diffusion coefficient can be related to the wall shear stress by the Chilton–Colburn analogy:

C2. Multiphase Flow and Flow Regime For multiphase flow the volume flow rates of the various phases affect the resulting flow pattern, called the “flow regime.” The flow regime is a primary consideration in the determination of flow-induced corrosion in multiphase flow. The basic criterion for flow-induced corrosion in multiphase

Superficial Liquid Velocity (Vsg), m/s

208


Pseudo Slug

Annular

Slug

Plug

Smooth Stratified

Wavy Stratified

Rolling Wave

Superficial Gas Velocity (Vsg), m/s

FIGURE 17.7. Typical gas–liquid flow regime map showing boundaries between various regions [11].

flow is that a liquid corrosive phase be present and be in contact with the metal surface. The flow regime affects the turbulence in the liquid phase, the location of the liquid phase in the flowing stream, and whether or not solids can settle out of the flowing stream. An example of a flow regime map showing the flow regimes encountered in multiphase gas–liquid horizontal flow is given in Figure 17.7 [11]. The bubbly flow regime occurs in liquid systems with low gas and high liquid rates. Stratified flow occurs at low liquid rates or when a liquid phase condenses. The liquid in stratified flow is generally stagnant or very low flow, with low turbulence.

BUBBLY

WAVY STRATIFIED

PLUG FLOW

PSEUDO SLUG

Slug flow occurs at higher liquid flow rates and is characterized by very high turbulence, with a corresponding influence on corrosion [10]. This flow regime has the most severe effect on corrosion. Annular flow occurs at very high gas rates over a wide range of liquid production rates. Turbulence in annular flow with respect to the effects on corrosion is not well characterized. A diagram of the flow regimes encountered in gas–water horizontal flow is given in Figure 17.8 [11] and hydrocarbon– water flow in Figure 17.9 [12]. When a hydrocarbon phase is present, the significant aspect of the flow regime with respect to corrosion is the separation of the hydrocarbon and water– liquid phases, or the level of turbulence required to eliminate the water phase as a separate entity contacting the metal wall. A diagram of the flow regimes encountered in gas– hydrocarbon–water horizontal flow and the flow regime map is given in Figure 17.10 [13]. Gas–hydrocarbon–water flow combines the characteristics of the gas–water and hydrocarbon–water flow regimes. The gas phase adds a high degree of turbulence, but the primary corrosion concern is the contact of water with the metal wall. The flow regime maps shown are for horizontal flow and are significantly altered for inclined flow [14]. The effect of inclination is particularly important for the occurrence and severity of slug flow with upward inclination, where both the frequency and intensity of the slug increases [14]. This is particularly significant due to the effect of slug flow on turbulence and corrosion [13]. A generalized summary of the types of corrosion expected for the various flow regimes, along with the location of free

SMOOTH STRATIFIED

ROLLING WAVE

SLUG FLOW

ANNULAR

FIGURE 17.8. Gas–water two-phase flow regimes in horizontal pipes [11].


SEGREGATED

SEMI-SEGREGATED

SEMI-MIXED

MIXED

CONCENTRIC

209

SLUG

DISPERSED

SEMI-DISPERSED

FIGURE 17.9. Hydrocarbon–water flow regimes in horizontal pipes [12].

BUBBLY

SMOOTH STRATIFIED

WAVY STRATIFIED

ROLLING WAVE

PLUG FLOW

SLUG FLOW

ANNULAR

PSEUDO SLUG

FIGURE 17.10. Gas–oil–water three-phase flow regimes in horizontal pipes [13].

water and the water turbulence, is outlined in Table 17.1 [15]. The flow regimes where either underdeposit corrosion or flow-induced corrosion is expected are given in this table.

. . . .

C3. Effect of Disturbed Flow Flow-induced corrosion failures in many environments occur in areas where fully developed flow patterns are disrupted, termed regions of disturbed flow [15–18]. Examples of these disturbed flow locations are:

Downstream Downstream Downstream Downstream inlets

of weld beads, at pipe joints and upsets of pipe fittings, at preexisting pits of valves, at bends and elbows in piping of orifice plates, at heat exchanger tube

The flow disturbance destroys the fully developed hydrodynamic boundary layer and the diffusion boundary layer. The effect of this boundary layer disruption is the production

210


TABLE 17.1. Flow Regime Influence on Corrosion Type [15] Flow Regime

Free Water Location

Water Turbulence

Corrosion Type

Gas/water Stratified flow

Bottom

Stagnant to laminar

Mostly bottom, mixed Circumferential

Very turbulent Turbulent

Underdeposit corrosion Pitting of stainless steels Flow-induced corrosion Flow-induced corrosion

Bottom, separated

Stagnant to laminar

Mostly bottom, mixed Circumferential

Very turbulent Possibly turbulent

Bottom

Stagnant to laminar

Mostly bottom, mixed Mixed

Laminar to turbulent Turbulent

Slug flow Annular flow Gas–liquid–water Stratified flow Slug flow Annular flow Liquid–water Segregated flow Mixed flow Dispersed flow

Underdeposit corrosion Pitting of stainless steels Flow-induced corrosion Flow-induced corrosion Underdeposit corrosion Pitting of stainless steels Underdeposit corrosion Flow-induced corrosiona

a

Only if the hydrocarbon water mixed phase is water external, that is, contacts the pipe wall.

of a steady-state condition as opposed to an equilibrium condition. In this case, the normal equilibrium corrosion reactions cannot be maintained, and a steady-state condition is established. The result is increased corrosion at the location of the flow disturbance, although the corrosion rate in the rest of the pipe wall is low. The corrosion pattern appears as large pits or corroded areas, often showing the signs of flow direction. A schematic of the effect of disturbed flow over a weld bead is shown in Figure 17.11 [15]. C4. Low-Turbulence-Flow Corrosion Low-turbulence-flow corrosion occurs in multiphase flow for a stratified flow regime and in single-phase flow containing solids where the turbulence is insufficient to keep the solids suspended. In low-turbulence conditions where solids are present deposits can accumulate at the bottom of a horizontal or slightly inclined line. A schematic showing the occurrence of corrosion under the accumulation of deposits in the bottom of a pipeline is shown in Figure 17.12. An example of the

severe corrosion that can occur in the bottom of a pipeline under deposits is shown in Figure 17.13. The presence of the solids deposit can also result in crevice corrosion of susceptible alloys, where the deposited solids form a crevice on the metal underneath. This can result in a differential aeration or differential metal ion concentration corrosion cell. This can result in a galvanic interaction between the steel under the deposit and that outside the deposit. The deposit can also allow bacteria to proliferate, resulting in microbiologically induced corrosion (MIC). Colonies of bacteria grow under the deposits and are protected by the deposits. For passive metals such as stainless steels, the area under the deposit can lose passivity, resulting in an active–passive galvanic couple. The area under the deposit can acidify just as in a traditional crevice, accelerating corrosion in that area. This is particularly severe when the water phase contains chlorides. Liquid and possibly solids accumulation in gas lines can be present in areas of liquid holdup. This occurs primarily on Equilibrium hydrodynamic and diffusion boundary layers reestablished

Equilibrium hydrodynamic and diffusion boundary layers

Equilibrium hydrodynamic and diffusion boundary layers destroyed, steady state established

STEEL

FIGURE 17.11. Effect of disrupted flow over weld bead on boundary layer [15].


Gas

Solids

Liquid Water

Pipe Wall

Under Deposit Corrosion

FIGURE 17.12. Occurrence of corrosion under water–solids deposits in bottom of pipeline at locations where liquid holdup occurs.

FIGURE 17.13. Corrosion in bottom of pipeline resulting from solids deposits under low-flow conditions.

Elevation Change

Liquid Water Solids

Pipeline Low Spot

Solids

Liquid Water

FIGURE 17.14. Locations where liquid holdup and solids deposition can occur.

211

212


areas of upward elevation change or at low spots, as shown in Figure 17.14. In these areas, gas moves across the top of the accumulated water, but the actual flow in the water is generally very slow to stagnant. Even with a high gas flow rate through the pipeline, the liquid flow in the areas of liquid holdup is sufficiently low to allow entrained solids to drop out from the liquid stream and settle to the bottom of the pipeline.

NOMENCLATURE C D d e i K L N P Rcor r T t U u x y yþ z

Concentration (mol/L) Diffusion coefficient (m2/s) Diameter, m Roughness factor, m Current density, A/m2 mass transfer coefficient (m/s) Characteristic length, m Mass flux density (mol/m2s) Pressure, Pa, kg/m s2 Corrosion rate, mm/y, mpy Radius or radial distance, m Temperature, C Time, s Mean velocity, m/s Local velocity fluctuation, m/s Direction parallel to surface, m Direction perpendicular to the surface, m Viscous length ¼ yut/n, dimensionless Direction parallel to the surface but perpendicular to the flow direction, m

Greek Symbols D d r m n t

Gradient of a property Boundary layer thickness (m) Density (kg/m3) Dynamic viscosity (kg/m s) Kinematic viscosity (m2/s) Shear stress (N/m2)

Subscripts b cor d f i j lim 0 t w

Boundary layer Corrosion Diffusion Final or ending Initial or beginning Species “j” Limiting Standard or primary Value at time t Wall or electrode surface

x y þ

Direction parallel to surface Direction perpendicular to surface Normalized dimensionless form

Dimensionless Groups ð2tÞ f Friction factor, ðrU 2 Þ ðrUdÞ Re Reynolds number, m m Sc Schmidt number, ðrdÞ ðkdÞ Sh Sherwood number, D REFERENCES 1. A. A. Townsend, The Structure of Turbulent Flow, Cambridge University Press, Cambridge, 1956, pp. 232–237. 2. H. Schlichting, Boundary-Layer Theory, McGraw-Hill, New York, 1979, p. 28. 3. S. J. Kline, W. C. Reynolds, F. A. Schraub, and P. W. Runstadler, “The Structure of Turbulent Boundary Layers,” J. Fluid Mech., 30, 741–773 (1967). 4. J. T. Davies, “Eddy Transfer Near Solid Surfaces,” in Turbulence Phenomena, Academic, New York, 1972, pp. 121–143. 5. M. T. Landahl and E. Mollo-Cristensen, Turbulence and Random Processes in Fluid Mechanics, Cambridge University Press, Cambridge, 1987, pp. 111–120. 6. A. K. Praturi and R. S. Brodkey, “A Stereoscopic Visual Study of Coherent Structures in Turbulent Shear Flow,” J. Fluid Mech., 89(Pt 2), 251–272 (1978). 7. G. R. Often and S. J. Kline, “A Proposed Model of the Bursting Process in Turbulent Boundary Layers,” J. Fluid Mech., 70, 209–228 (1975). 8. S. K. Robinson, “Coherent Motions in the Turbulent Boundary Layer,” Annu. Rev. Fluid Mech., 23, 601–639 (1991). 9. A. S. Thomas and M. K. Bull, “On the Role of Wall Pressure Fluctuations in Deterministic Motions in the Turbulent Boundary Layer,” J. Fluid Mech., 128, 283–322 (1983). 10. N. Ibl and O. Dossenbach, “Convective Mass Transport,” in Comprehensive Treatise of Electrochemistry, Vol. 6, Electrodics: Transport, E. Yeager, J. O’M. Bockris, B. E. Conway, and S. Sarangapani (Eds.), Plenum, New York, 1983, pp. 133–237. 11. A. H. Lee, J. Y. Sun, and W. P. Jepson, “Study of Flow Regime Transitions of Oil-Water-Gas Mixtures in Horizontal Pipelines,” 3rd International Conference ISOPE, Vol. II, 1993, pp. 159–164. 12. A. Maholtra,“A Study of Oil/Water Flow Characteristics in Horizontal Pipes,” MS Thesis, Ohio University, Athens, OH, Oct., 1996. 13. X. Zhou and W. P. Jepson, “Corrosion in Three-Phase Oil/Water/Gas Slug Flow in Horizontal Pipes,” Paper No.

REFERENCES

94026, CORROSION/94, NACE International, New Orleans, LA, Mar., 1994. 14. C. Kang, R. Wilkins, and W. P. Jepson, “The Effect of Slug Frequency on Corrosion in High Pressure Inclined Pipelines,” Paper No. 96020, CORROSION/96, NACE International, Denver, CO, Mar. 1996. 15. K. D. Efird, “Disturbed Flow and Flow Accelerated Corrosion in Oil and Gas Production,” Proceedings: ASME Energy Resources Technology Conference, Houston, TX, Feb. 1998.

213

16. J. Postlethwaite, S. Nesic, G. Adamopoulos, and D. J. Bergstrom, “Predictive Modeling for Erosion Corrosion under Disturbed Flow Conditions,” Corros. Sci., 35, 627–633 (1993). 17. S. Nesic and J. Postlethwaite, “Relationship between the Structure of Disturbed Flow and Erosion Corrosion,” Corrosion, 46, 874–880 (1990). 18. G. Schmitt and T. Gudde, “Local Mass Transport Coefficients and Wall Shear Stresses at Flow Disturbances,” Paper No. 95102, CORROSION/95, Orlando, FL, Mar. 1995.

18 EROSION–CORROSION IN SINGLE- AND MULTIPHASE FLOW J. POSTLETHWAITE* Department of Chemical Engineering, University of Saskatchewan, Saskatoon, Saskatchewan, Canada S. NESI c

Institute for Corrosion and Multiphase Flow Technology, Ohio University, Athens, Ohio

A. B. C. D. E. F.

Introduction Flow conditions Protective films Erosion–corrosion rate Relative roles of erosion and corrosion Erosion–corrosion mechanisms F1. Turbulent flow F2. Solid-particle impingement F3. Liquid-droplet impingement F4. Air bubble impingement F5. Cavitation F6. Wire drawing F7. Flow-enhanced film dissolution References

A. INTRODUCTION Erosion–corrosion encompasses a wide range of flowinduced corrosion processes. Flowing fluids can damage protective films on metals resulting in greatly accelerated corrosion. Damage to the films may be the result of mechanical forces or flow-enhanced dissolution and the accelerated corrosion may be accompanied by erosion of the underlying

*

Retired.

metal. This conjoint action of erosion and corrosion is known as erosion–corrosion [1]. Impingement attack by liquid droplets and solid particles and cavitation attack are also included here under this broad definition of erosion– corrosion. In practice, the relative contributions of accelerated corrosion and erosion to the total metal loss vary widely with the type of erosion–corrosion and the hydrodynamic intensity of the flow.

B. FLOW CONDITIONS Erosion–corrosion normally occurs under turbulent-flow conditions. The flowing fluid may be single phase (Fig. 18.1) as in the erosion–corrosion of copper tubing by potable water (Fig. 18.2). Multiphase flows [2–4] (Fig. 18.3) with various combinations of gas, water, oil, and sand can cause severe erosion–corrosion of oil/gas production systems [5–7] as shown in Figures 18.4–18.6 The most severe erosion–corrosion problems occur under conditions of disturbed turbulent flow [8, 9] at sudden changes in the flow system geometry, such as bends, heat exchanger tube inlets, orifice plates, valves, fittings, and turbomachinery, including pumps, compressors, turbines, and propellers. Surface defects in the form of small protrusions or depressions such as corrosion pits, deposits, and weld beads can give rise to disturbed flow on a smaller scale but sufficient to initiate erosion–corrosion [10, 11]. The presence of suspended solid particles, gas bubbles, or vapor bubbles in aqueous flow and liquid droplets in high-speed gas flow can be especially damaging.


215

216

EROSION–CORROSION IN SINGLE- AND MULTIPHASE FLOW

Diffusion barriers are typically formed by anodic dissolution followed by precipitation whereas passive films are formed by direct oxidation without the metal ions entering the solution. The softer thicker diffusion barriers are more easily damaged and more slowly repaired than the passive films, and passive alloys can withstand much more severe service conditions.

D. EROSION–CORROSION RATE FIGURE 18.1. Single-phase pipe flow. (a) Developed laminar flow, showing parabolic velocity profile. (b) Developed turbulent flow, showing logarithmic velocity profile with large gradients near the wall (nondisturbed flow). (c) Disturbed turbulent flow with separation, recirculation, and reattachment, showing complex velocity field. (Adapted from Heitz [18].)

C. PROTECTIVE FILMS Most metals and alloys used in industry owe their corrosion resistance to the formation and retention of a protective film. Protective films fall into two categories: .

.

The relatively thick porous diffusion barriers such as those formed on carbon steel (red rust) and copper (cuprous oxide) The thin invisible passive films such as those on stainless steels, nickel alloys, and other passive metals such as titanium

The effect of film destruction on the corrosion rate is illustrated by the following example. Carbon steel pipe carrying water is usually protected by a film of rust, which slows down the rate of mass transfer of dissolved oxygen to the pipe wall. The resulting corrosion rates are typically 10 mm/y. Remedy: Continuous inhibitor injection. (From Houghton and Westermark [6]. Reprinted by permission of NACE International.)

FIGURE 18.5. Impingement-type erosion–corrosion of AISI 4140 (UNS 941400) 115-mm flow coupling and subsurface safety valve in natural gas condensate production. Minor species: CO2 and H2O. Temperature: 79 C. Exit velocity from valve 91 m/s. Remedy: replacement with 13 Cr martensitic ss. (From Houghton and Westermark [6]. Reprinted by permission of NACE International.)

protective film can be estimated (Table 18.1) on the basis of the exponent y in the relationship between the metal loss rate and the bulk flow velocity, ub [2]: erosion þ corrosion / uyb The value of y will depend on the relative contributions of corrosion and erosion to the total metal loss. In disturbed flow it is the flow characteristics in the direct vicinity of the wall rather than the bulk flow velocity that are important. In practice, however, the superficial flow velocity is the flow parameter that is readily measured and controlled [2].

F. EROSION–CORROSION MECHANISMS The sources of the various mechanical forces involved in the erosion of protective films and/or the underlying metal at flow system walls (Fig. 18.8) are: . . .

Turbulent flow, fluctuating shear stresses, and pressure impacts Impact of suspended solid particles Impact of suspended liquid droplets in high-speed gas flow

EROSION–CORROSION MECHANISMS

219

FIGURE 18.6. Erosion of tungsten carbide choke beans inside a steel holder from an oil well with sand production. Note the highly polished, streamlined appearance of the erosion pattern. Subcritical flow occurred in this well and these assemblies lasted between 6 and 9 months. (From Smart [7]. Reprinted by permission of NACE International.)

TABLE 18.1. Flow Velocity as Diagnostic Tool for Erosion– Corrosion Rates Following Damage to Protective Film Mechanism of Metal Loss Corrosion Liquid-phase mass transfer control Charge transfer (activation) control Mixed (charge/mass transfer) control Activation/repassivation (passive films) Erosion Solid-particle impingement Liquid-droplet impingement in highspeed gas flow Cavitation attack

. .

Velocity Exponent, y 0.8–1 0 0–1 1 2–3 5–8 5–8

Impact of suspended gas bubbles in aqueous flow Violent collapse of vapor bubbles following cavitation

Flow-enhanced film dissolution and thinning are “chemical” forms of protective film erosion leading to accelerated corrosion of the underlying metal. F1. Turbulent Flow FIGURE 18.7. Effect of pipe wall roughness on mass transfer rates. (a) Rough surface, preroughened by erosion–corrosion prior to mass transfer study.(b) Roughness profile details: roughness height e 0.2 mm, pitch/height in the range 5–10. (c) The ratio of measured mass transfer coefficients for preroughened and smooth surfaces at pipe wall. (Adapted from Postlethwaite and Lotz [17].)

Erosion–corrosion occurs in single-phase turbulent pipe flow. The exact mechanism of protective film damage is still in doubt. There is uncertainty regarding the roles of mechanical forces and mass transfer in film disruption [20–22] since both are directly related to the turbulence intensity.

220

EROSION–CORROSION IN SINGLE- AND MULTIPHASE FLOW

flow mass transfer

Single phase heat transfer

fluctuating pressure

fluctuating shear stress

protective film

wall Liquid/solid particles

Liquid/air bubbles

Gas/liquid droplets

Liquid/vapor bubbles violent collapse of vapor bubbles formed upstream

FIGURE 18.8. Interaction of flowing fluid with wall of the flow system leading to erosion–corrosion. (Adapted from Heitz [18].)

Distinct “breakaway” velocities above which damage to normally protective films occurs in copper alloy tubes are observed [20] and this gave rise to the concept of a critical shear stress for film disruption. The wall shear stress tw for flow in pipes is given by tw ¼ f ð12r u2b Þ where f is the Fanning friction factor [23]. In general [24] e e0 f ¼ f Re; ; ; m d d where e is a measure of the size of the roughness projections, e0 is a measure of the arrangement of the roughness elements, and m is a form factor that is dependent on the shape of roughness elements. For flow in smooth pipes f ¼ 0:046 Re 0:2 is valid to higher Reynolds numbers than the Blasius equation [25]. Values of the Fanning friction factor for both smooth and rough commercial pipes and tubes can be obtained from a Moody* chart [23]. For fully rough† pipe flow, f is independent of Re, and tw / u2b .

*

The Fanning friction factor should not be confused with the Darcy friction factor, which is four times greater and is used on some Moody charts. † In flow through rough pipes, the roughness elements penetrate the boundary layer into the main fluid stream at high Re resulting in form drag, and, under fully rough conditions, the viscous forces become negligible.

Early attempts [26] to correlate film damage with the wall shear stresses in pipe flow based on the bulk flow velocity resulted in shear stress values (Table 18.2) that seem too low [20] to strip a corrosion product film unless it is very loosely adhering. The “breakaway” velocities ucr shown in Table 18.2 were calculated from critical shear stress values. Design velocities for heat exchanger tubing should not be based on critical shear stress values for fully developed non disturbed tube flow since the flow is both developing and disturbed at the tube inlets. In fact, disturbed flow must be allowed for in most industrial systems. The values calculated from 0.5ucr are also shown and can be compared to the “maximum recommended design velocities” [27]. The design velocity must clearly be chosen with care with previous experience in similar systems being the best guide. In turbulent flow there are fluctuating shear stresses and pressures at the wall. The largest values are obtained during quasi-cyclic bursting events close to the wall which are said to be responsible for most of the turbulent energy production in the entire wall-bounded shear flow [28, 29]. This is true of both nondisturbed and disturbed pipe flows. In practice, film removal in single-phase aqueous flow is invariably associated with the vortices created under disturbed flow conditions produced by sudden macro- or microscale changes in the flow geometry. It is quite difficult to avoid microscale geometric imperfections in a flow system during experimental studies to determine critical shear stress values. The films on passive alloys, such as stainless steels, are not usually damaged by single-phase aqueous flow [18, 30]. Rust films on carbon steel at ambient temperature are more mechanically stable than the protective films formed on copper. The compactness and protectiveness of rust films (FeOOH) increase with velocity [31]. The magnetite films formed in high-temperature water and sulfide films formed

EROSION–CORROSION MECHANISMS

TABLE 18.2.

221

Critical Flow Parameters for Copper Alloy Tubing in Seawater

Alloy Cupro nickel with Cr 70–30 Cupro nickel 90–10 Cupro nickel Aluminum bronzes Arsenical Al brass Inhib. Admiralty Low Si bronze P deoxidized copper

Critical Shear Stressa (N/m2)

Critical Velocity 25-mm tubeb (m/s)

297 48 43 — 19 — — 9.6

12.6 4.6 4.3 — 2.7 — — 1.9

“Design” Velocityc Based on 50% tw,crit (m/s) 8.6 3.1 2.9 — 1.9 — — 1.3

Accepted Maximum Design Velocityd (m/s) 9 4.5–4.6 3–3.6 2.7 2.4 1.2–1.8 0.9 0.6–0.9

a

From Efird [26], rectangular duct. Critical velocities, calculated using Efird shear stress values. c Calculated values based on 50% of critical shear stress. d “Accepted” maximum tubular design velocities [27]. b

on carbon steel have a high mechanical stability in singlephase aqueous flow [18]. Thus the velocity limit, around 1 m/s, suggested for copper tubing [32, 33] is lower than the velocities tolerated by carbon steel and low-alloy piping at ambient temperatures and much lower than the allowable velocities at elevated temperatures. In slug flow [Fig. 18.3(c)] large but short-lived shear stress fluctuations are observed [2]. The slugs of liquid travel at much higher velocities than is encountered in singlephase liquid flows and have been suspected to be the cause of severe erosion–corrosion problems in oil/gas production systems. F2. Solid-Particle Impingement Solid-particle impacts can damage both types of protective films (thick diffusion barriers and thin passive films) leading to erosion–corrosion. The particles may also erode the underlying metal, adding to the overall metal loss. As might be expected, the erosion rate ER is a function of the kinetic energy of the particles (proportional to u2p ) and the frequency of impacts (proportional to up), and to a first approximation [2], ER / u3p The kinetic energy of the particle normal to the wall will be determined by the impact velocity and impact angle along with the particle density and size. The impacts of small particles in aqueous slurries may be damped by the presence of a thick boundary layer at the wall [36]. In practice, hard abrasive slurries are often transported through mildsteel pipes with the particles finely ground. The erosion is reduced because of the smaller particle size and the lower flow velocity required to maintain the solids in suspension [37].

Both the impact velocity and the impact angle will be strongly affected by disturbed flow and this is where the most severe erosion is found in practice. Impact angles in nondisturbed turbulent pipe flow are 50 C) and lower pH ( 0) and h is the metal recession (increase in oxide film thickness dx ¼ PBR dh). Equation (20.45) allows the tangential stresses stOx to be calculated assuming linear elastic behavior. The magnitude of the maximum radial stresses, srOx , is given by srOx ¼

PBR h t sOx Rs

ð20:46Þ

The relationship between the signs of the tangential and radial strains and stresses in the scale and at the metal–oxide interface, respectively, are sign etOx ¼ signð erOx Rs Þ

ð20:47Þ

signð srOx Rs Þ

ð20:48Þ

sign stOx

¼

A plus sign indicates tensile stress, a minus sign means compressive stress. The sign and level of the stresses in the scale depend on its growth direction and on the radius of service curvature as well as the PBR. The latter was introduced in 1925 in order to explain the formation of growth stresses during oxidation and describes the volume change that is involved in the transition from the metal lattice to the TABLE 20.1.

cation lattice of the oxide when only the oxygen anions are diffusing. In other words, the PBR corresponds to the ratio of the volume per metal ion in the oxide to the volume per metal atom in the metal [13]: PBR ¼

ð20:44Þ

where a is the fraction of oxide formed on the scale surface, 1 a is the fraction of oxide formed at the metal–oxide interface, V is the volume fraction of metal consumed in the oxidation by injecting vacancies into the metal, 1 V is the volume fraction of metal consumed in the oxidation which originates directly from the metal surface, and PBR is the Pilling–Bedworth ratio (see later in this section). Oxidation leads to an increase in the strain in the circumferential direction (tangential strain etOx ) with a rate of

259

1 Veq ðMa Ob Þ a Veq ðMÞ

ð20:49Þ

It was argued that if the PBR was less than 1, the growth stresses would be tensile and the oxide would crack and not cover the entire metal surface. As indicated in Table 20.1, alkali and alkaline earth metals belong to this class of materials. On the other hand, if the PBR was higher than 1, compressive stresses would develop and the oxide could be protective, at least during the early stages of oxidation. The majority of metals fall into this category. We now know that the Pilling–Bedworth rule regarding protective behavior exhibits several exceptions. Important examples are tantalum or niobium where, even though the PBR is substantially larger than unity, cracks develop in the oxide scale after extended exposure and these produce nonprotective conditions. Whereas the Pilling–Bedworth paper was a significant advance at the time, it is now recognized that the approach was incomplete and that the influence of the difference between the molar volume of metal and oxide depends on the oxide growth mechanism. However, the Pilling–Bedworth approach may be of great help for the assessment of the geometrically induced growth stresses, as shown earlier. D3.2. Thermal Stresses. In most applications, hightemperature alloys are subjected to temperature fluctuations even under nominally isothermal conditions. In this case, though, the resultant stresses in the oxide layer resulting from the difference in the coefficient of thermal expansion (CTE) of the metal and oxide (see Table 20.2) are small and may be neglected. This will not be the case, however, for large thermal cycles or during cooling to room temperature when large stresses, perhaps of 1-GPa order, are produced in the oxide layer. Metals have generally a higher coefficient of thermal expansion than oxides (Table 20.2), and consequently, tensile stresses are induced in the oxide scale on heating and compressive stresses during cooling. The thermally induced stresses can be calculated from the coefficients of thermal expansion according to the equation [94]

Pilling–Bedworth Ratios for Some Metal–Oxygen Systems

Oxide PBR

K2O 0.45

CaO 0.64

MgO 0.8

CeO2 0.90

Na2O 0.97

CdO 1.21

Al2O3 1.28

ZnO 1.55

ZrO2 1.56

Oxide PBR

Cu2O 1.64

NiO 1.65

FeO 1.7

TiO2 1.73

CoO 1.86

SiO2 1.9

Cr2O3 2.07

Ta2O5 2.5

Nb2O5 2.7

260

HIGH-TEMPERATURE OXIDATION

TABLE 20.2. Linear Coefficients of Thermal Expansion of Metals and Oxides, C 1 [95] System Fe/FeO Fe/Fe2O3 Ni/NiO Co/CoO Cr/Cr2O3 Cu/Cu2O Cu/CuO

Oxide: 106aox

Metal: 106aM

Ratio:aM =aOX

12.2 14.9 17.1 15.0 7.3 4.3 9.3

15.3 15.3 17.3 14.0 9.5 18.6 18.6

1.25 1.03 1.03 0.93 1.30 4.32 2.0

EOx DTðaM aOx Þ stherm ¼ EOx dOx ð1 nÞ 1þ2 EM d M

ð20:50Þ

where a is the CTE for the metal and the oxide, E is Young’s modulus for the metal and the oxide, d is the thickness for the metal and the oxide, respectively, and n is Poisson’s ratio. Here, DT stands for the temperature change. The CTEs for technical materials can be found in many of the materials producers’ brochures, and those for corrosion products are given in the literature [95]. In most cases, the CTEs can be approximated by linear behavior in the temperature range concerned, but in some cases, where phase changes occur in the scale during the temperature change, nonlinear temperature dependence is found for the CTE. This is, for example, the case for several sulfide layers [96] and is particularly important for magnetite and some iron-based spinels [97], affecting stresses in oxide scales on low-alloy steels. This naturally decisively affects the stress situation in the oxide scales on low-alloy steels [98]. In the temperature range between about 600 and 450 C, the magnetite partial layer is under tensile stress when cooling from 600 C. At lower temperatures, this oxide partial layer may come under compressive stresses, depending on the metallic substrate and its CTE. In the hematite layer, the stresses are always compressive, as the CTE always lies below that of the low-alloy steel (the exception is 9% chromium steel at temperatures below 150 C). D3.3. Mechanical Scale Failure. Growth stresses and thermal stresses may be relieved through various mechanisms that could operate simultaneously: . . . .

Plastic deformation of the oxide scale Plastic deformation of the metal substrate Spalling of the oxide from the alloy Cracking of the scale

When plastic deformation is not sufficient for stress relief, cracking may develop in the scale. It is the more efficient

relaxation mechanism but will result in a sudden increase in corrosion rate. The metal oxidation may exhibit repeated regular sequences of cracking and healing of the scale. Under tensile stresses (heating to temperatures higher than the oxidation temperature or over convex regions of a nonplanar surface) cracks appear as soon as the elastic fracture strain is reached. This critical value will be significantly less than 1% even at high temperatures. Under compressive stresses, the degradation leads to spallation and the mechanisms are more complex. Two processes are necessary to produce spalling: transverse cracking through the oxide and decohesion along the metal–oxide interface. Two routes of spallation have been identified: the case corresponding to a low cohesive strength of the oxide and a high adhesive strength of the scale on the substrate surface (route 1: cracking of the oxide before decohesion) and the case corresponding to a high cohesive strength of the oxide and poor adhesion of the oxide to the metal (route 2: decohesion before metal cracking). Figure 20.9 illustrates these two distinct mechanisms.

E. COMPACT SCALE GROWTH E1. Elementary Chemical Steps The overall oxidation reaction of a metal M may be written as b aM þ O2 ¼ Ma Ob 2

ð20:51Þ

The reaction can proceed only if diffusion of matter (oxygen or metal) occurs through the solid scale MaOb. If the scale is porous, mass transport occurs by oxygen diffusion; if the scale is compact, mass transport occurs by means of solidstate diffusion. In the latter case, the oxidation mechanism consists of at least four steps [100–104]: . . . .

Surface step—oxygen adsorption on the oxide External step—matter exchange at the adsorbed phase/ oxide interphase Diffusion step—ionic transport through the oxide scale Internal step—matter exchange at the oxide/metal interphase

Of course, mass transport by migration of ionized point defects is accompanied by simultaneous electrical transport, which complicates the process. In this section, the kinetics of the oxidation process leading to the growth of a compact scale are analyzed. Let a chemical elementary step be a local reversible reaction that occurs without the formation of a distinct intermediate product, that is, the reaction proceeds in a single step, and let an interphase elementary step be a chemical

COMPACT SCALE GROWTH

261

FIGURE 20.9. Cracking and spallation by compressive oxide stresses [99].

process involving matter exchange between two distinct phases. The main difficulty encountered for the formulation of these steps relates to the conditions required for the application of classical theories (Eyring theory) for the calculation of the step rates [105]. We will have to choose, for the adsorbed phase as well as for the oxide, structural models that exhibit ideal behavior for the reacting species in their own phase. Moreover, we will assume that the theory of absolute rates can, under these conditions, be extended to heterogeneous elementary steps involving matter exchange between two different phases. On a solid surface, the atomic environment is modified in comparison with that in the bulk. The resulting imbalance of the forces in the surface of solids produces attractive forces for gas molecules or atoms. The phenomenon of adsorption can then produce an excess of gas atoms or molecules on the surface compared with the concentration in the adjacent gas phase. For adsorption to occur spontaneously, the process must produce a decrease in free energy, but since that is also a decrease in system entropy, adsorption is always an exothermic process. Consequently, the amount of adsorbed gas at equilibrium at constant pressure (the “adsorption isobar”) decreases with increasing temperature. Depending on the nature of the forces involved, adsorption processes may be classified as physical adsorption (also termed van der Waals adsorption or physisorption) or chemical adsorption, usually abbreviated to chemisorption [106,

107]. Physisorption is generally quasi-instantaneous while chemisorption often proceeds slowly, involving an activation energy Ea. Thus, the chemical adsorption rate becomes appreciable only at sufficiently high temperature. Many theories and models have been proposed to explain the shape of adsorption isotherms that represent the variation of adsorbed volume as a function of gas pressure or of the p/p0 ratio (p0 is the saturation vapor pressure at the experimental temperature). The description of monolayer adsorption can be made using as variable the fraction of the available adsorption sites which are occupied by adsorbed atoms or molecules, u ¼ s=s0 , where s0 is the number of adsorption sites that are initially available per unit surface area and s is the number of occupied surface sites per surface area unit (thus, u is the fraction of occupied sites). Chemisorption involves partial electronic transfer between adsorbed molecules and the substrate. The solid surface appears inhomogeneous and exhibits specific “active” sites on which chemisorption takes place preferentially. When temperature is increased, the amount of adsorbed gas by chemisorption increases because it is an activated process; then the adsorption isobar passes through a maximum because chemisorption is an exothermic process. Since the establishment of high-temperature oxidation requires at least the presence of one monolayer on the surface, it can be assumed that a chemisorbed phase is produced as a surface step.

262


In this discussion, this adsorbed phase is considered as a two-dimensional solution (i.e., sorption of a monolayer) of free surface sites, s, and of occupied sites leading to the formation of a superficial compound referred to as O-s (atomically chemisorbed oxygen) where O is a particle of the gas phase (O2). Such a solution may be considered as ideal since it is assumed that no interaction occurs between the free and occupied sites. The sorption process may then be described by an equation representing the balance between two opposite reactions with rate constants Ka0 and K 00a , respectively: O2 þ 2s ¼ 2O-s

ð20:52Þ

with Ka ¼ Ka0 =Ka00 denoting the equilibrium constant of the elementary step of sorption. Under these conditions, if p denotes the partial pressure of the pO2 species, ueq the fractional surface coverage, the law of mass action applied to the equilibrium leads to

Then, in the case of an n-type semiconductor with metal vacancies, the external step can be described by 00 þ 2h þ s O-s ¼ OO þ VM .

with Ke ¼ Ke0 =Ke00 denoting the equilibrium constant of the external step of matter exchange. The external reaction rate is given by dne . 00 ¼ Ke0 Q Ke00 ½VM e ½h 2e ð1 QÞ dt

h 00 i dne 00 eq 3 ¼ Ke0 Q 1 ½VM e =½ V M e dt

ð20:53Þ

dnO-s ¼ Ka0 ð1 QÞ2 pO2 Ka00 Q2 dt

ð20:54Þ

"

1Q Q

#

2 Ka pO2 1

ð20:55Þ

with Ka pO2 ¼

Qeq 1 Qeq

2 ð20:56Þ

The proposed model to account for the external and internal steps uses the concept of point defects and, for simplicity, it can be assumed that only one defect is predominant in the 00 or the oxygen lattice, that is, either the metal vacancy VM 00 interstitial O i for the case of a p-type semiconductor and either the oxygen vacancy VO.. or the metal interstitial Mi.. for the case of an n-type semiconductor. The defects are then generated at the external interphase (cationic diffusion, p-type semiconductor) or at the internal interphase (anionic diffusion, n-type semiconductor) (see Fig. 20.6).

ð20:60Þ

00 VM þ 2h þ M ¼ MM

ð20:61Þ

dni 00 3 00 eq3 ¼ 4Ki0 ½VM i ½V M i dt

ð20:62Þ

.

with eq

dnO-s Q2 ¼ Ka0 dt Ka

ð1=4ÞKe Q 1Q

Again for a p-type semiconductor with metal vacancies, we have for the internal step

00 ½V M i ¼

or

ð20:59Þ

with eq

indicating that the sorption is governed by the Langmuir equation (dissociation occurs on chemisorption, which is likely at high temperature). If dnO-s is the number of O-s particles formed by adsorption per unit area during the time dt, the rate of the sorption process can be given by the equations

ð20:58Þ

Under equilibrium conditions, for a given value of u, we have

00 ½V M e ¼

pffiffiffiffiffiffiffiffiffiffiffiffi Ka pO2 eq pffiffiffiffiffiffiffiffiffiffiffiffi Q ¼ 1 þ Ka pO2

ð20:57Þ

1 1 1=3 4 Ki

ð20:63Þ

In a thick, growing oxide scale, a positive space charge can appear in contact with the metal balanced by a negative space charge localized near the external interphase. This charge distribution induces, at any point within the scale, an electric field that both accelerates the positively charged ionic defects and slows down the negatively charged electronic defects until no net electrical current flows through the scale. Consequently, stationary concentration profiles are established in the MO scale, the electrically neutral zone extending over practically all of the scale thickness [108, 109]. Taking into account the possible electrical potential gradient through the scale, the particle fluxes are given by

qcd qf Cd Ud Jz ¼ Dd qx qx Je ¼ De

qce qf Ce Ue qx qx

ð20:64Þ

ð20:65Þ

263

COMPACT SCALE GROWTH

In these expressions, Cd (Ce) is the volume concentration of point (electronic) defects with z degree of ionization, Dd (De) is the diffusion coefficient, Ud (Ue) is the electric mobility of these charge carriers, and f the electric potential at the coordinate x. The Cd and [d] are linked by the relationship Cd ¼ [d]/W, W being the volume of 1 mol of oxide. A positive sign has to be used if the defects are positively charged and a negative sign for negatively charged defects. A brief kinetic analysis of the four elementary steps for reaction (51) has been presented in this section. The rate expressions of the elementary steps were expressed in a form involving the deviation from equilibrium, that is, within the framework of the thermodynamics of irreversible processes. The general system of equations relating to the growth of an oxide scale MO can then be established by expressing the mass balance at both sides of each interphase in the adsorbed phase and within the scale. These equations, being differential, cannot be solved analytically. However, it is possible to make the simplifying assumption that the concentrations tend to become time independent, that is, a quasi–steady state develops. On this basis, the system of differential equations allowing calculation of the reaction rate can be solved analytically, leading to considerable simplifications. It should be recalled that if one of the rate constants or the diffusion coefficient has a finite value, all the other constants having very large values, we deal with what is called a pure regime [110–113]. In all other cases, it is called a mixed regime [75, 114–116]. Hereafter, we describe pure diffusional regimes.

where

00 00 00 0 ½V M : K 0 ¼ DV M00 W ½VM

ð20:68Þ

Integrating and noting that y ¼ 0 at t ¼ 0 we obtain y2 ¼ 2K 0 t

ð20:69Þ

which is the common parabolic rate law. Furthermore, it has been shown that the cation vacancy concentration is related to the oxygen partial pressure by the equation 00 ½V M ¼ const pO2

1=n

ð20:70Þ

The variation of the parabolic rate constant with oxygen partial pressure can be predicted by h i 1=n ðp0O2 Þ1=n K 0 a p00O2

ð20:71Þ

and since p0O2 is usually negligible compared with pO00 2 , we have 1=n K 0 a p00O2 ð20:72Þ Clearly, the concentration gradient in the scale never equals zero and, therefore, scale growth never stops. Originally, Eq. (20.69) was derived by Wagner in a theoretical, detailed analysis of the electrochemical potential situation and the transport conditions in the scale. Figure 20.10 gives a summary of the conditions for which the theory is valid. Assumptions are as follows:

E2. Diffusion-Controlled Oxidation Assuming that cationic transport across the growing oxide layer controls the rate of scaling and that thermodynamic equilibrium is established at each interphase, the outward cation flux JM2 þ is equal and opposite to the inward flux of cation vacancies and we can write

. . . .

JM2 þ ¼ JV M00 ¼ DV M00

00 00 00 0 ½VM ½V M y

ð20:66Þ .

where y is the oxide thickness, DV M00 is the diffusion coefficient 00 0 00 00 and ½VM are the vacancy for cation vacancies, and ½VM concentrations at the scale–metal and scale–gas interphases, respectively. Since there is thermodynamic equilibrium at each in00 00 00 0 ½V M is constant and we terphase, the volume of ½VM have J M2 þ ¼

1 dy 1 K 0 ¼ W dt W y

ð20:67Þ

.

.

The oxide layer is a compact, perfectly adherent scale. Migration of ions or electrons across the scale is the rate-controlling process. Thermodynamic equilibrium is established at both the metal–scale and scale–gas interfaces. The oxide scale shows only small deviations from stoichiometry and, hence, the ionic fluxes are independent of position within the scale. Thermodynamic equilibrium is established locally throughout the scale. The scale is thick compared with the distances over which space charge effects (electrical double layer) occur. Oxygen solubility in the metal may be neglected.

This model led to the final equation of the parabolic rate constant, which is [11, 109, 117, 118] Kp0 ¼

1 2

ð p00O 2

pO0 2

zc DM þ DO d ln pO2 jza j

ð20:73Þ

264


FIGURE 20.10. Diagram of scale forming according to Wagner’s model [12].

where Kp0 , DM, and DO are given in units of centimeters squared per second, with DM and DO the self-diffusion coefficients for random diffusion of the respective ions (metal and oxygen) and z the valence of the respective ion (anion and cation). The parabolic rate law that was derived for thickness growth can also be modified for weight gain by oxidation. In this case, Kp0 has to be replaced by Kp00 according to the equation Kp00

¼2

ð2

Kp0

ð20:74Þ

z

Ð Values of z are given in Table 20.3 for several oxides. These values are, however, based on the assumption that the scale is free of pores and cavities and consists of only one phase. Under practical conditions, this is not usually the case, and therefore these values should only be used as estimations. Ð TABLE 20.3. Calculation Factor z for Conversion of Mass Gain Data into Scale Thickness Data Using Equation (20.74) for Several Oxides and Sulfides [119] Ð 3 Corrosion Product z ðg=cm Þ FeO Fe3O4 Fe2O3 Cr2O3 FeCr2O4 FeS FeS2

1.28 1.43 1.57 1.64 1.45 1.76 2.60

The oxidation rate constant Kp is the most important parameter for describing oxidation resistance. If Kp is low, the overall oxidation rate is low and metal consumption occurs at a very low rate. This is typical for protective oxidation. If Kp is high, metal consumption occurs at a high rate and the case of nonprotective oxidation exists. From Kp the metal consumption rate can also be calculated [120]. This requires knowledge of the stoichiometry of the oxide and the specific weight values as well as the molar weights of the reactants. Then the metal consumption can be calculated. The real value of Wagner’s analysis lies in providing a complete mechanistic understanding of the process of high-temperature oxidation under the conditions set out. The predictions of Wagner’s theory for n-type and p-type oxides have been extensively examined by several workers [120– 129]. For many systems, the obtained rate constants are generally several orders of magnitude larger than those which one would calculate from lattice diffusion data from Eq. (20.73). This discrepancy indicates that “short-circuit transport” is contributing to growth of the oxide film. E3. Short-Circuit Diffusion Lattice diffusion is the dominant process of mass transport at high temperature provided there is a sufficiently high defect concentration, but mass transport can also occur along dislocations or grain boundaries. Thus, the overall transport in polycrystalline oxide scales generally results from two fluxes of matter in parallel: an intragranular flux JV and a flux along grain boundaries Jgb.

265

MULTILAYERED SCALE GROWTH

The oxidation rate may be expressed as a function of the effective diffusion coefficient defined by Eq. (20.43) (see Section D2). The scale growth rate obeys the differential equation dy DC ¼ WDeff dt y

ð20:75Þ

M

Y1

Y2

M2O

MO

MM1

MM2

e´

e´

VM1

VM2

O2

where DC is the defect concentration difference through the oxide scale and W the molar volume of the oxide. Integration of Eq. (20.75) leads to y2 ¼ 2Deff DC t ¼ K 0pðeffÞ t

ð20:76Þ

where K 0 pðeffÞ is the effective parabolic constant. Assuming that the oxide is dense and pure and the grains are spherical, the following rate equation can be obtained [110, 130–132]:

1 y2 ¼ 2W DC DV t þ 4Dgb ðKt þ r20 Þ1=2 K

Po´2

JM ¼

F1. Compact Subscales The theory of multilayered scale growth on pure metals has been treated by Yurek et al. [133]. The hypothetical system treated is shown in Figure 20.11. It is assumed that the growth of both scales is diffusion controlled with the outward migration of cations large relative to the inward migration of anions. The flux of cations in each oxide is assumed to be independent of distance. Each oxide exhibits predominantly electronic conductivity and local equilibrium exists at the phase boundaries. The total oxidation reaction is ð20:78Þ

The cation vacancies are assumed to be neutral. The cation flux in subscale 1 is

1 Kp0 1 W 1 y1

ð20:79Þ

and the amounts Q of metal consumed per unit area to form layers 1 and 2 are, respectively, Q1 ¼

2y1 W2

Q2 ¼

y2 W2

ð20:80Þ

The fractions of the cation flux involved in the growth of subscales (1) and (2) are, respectively,

F. MULTILAYERED SCALE GROWTH The formation of several corrosion scales on metals or alloys is often observed. In order to understand clearly the growth mechanisms of these scales, it is of major importance to know whether they form concurrently or sequentially.

Po´´ 2

FIGURE 20.11. Schematic diagram of hypothetical two-layered scale.

ð20:77Þ

where r0 is the initial grain radius and K is a constant which is a function of, among others, the surface energy of grain boundaries, their thickness, and the diffusion coefficient of matter across the grain boundaries. Further numerical development of this simple model has allowed a grain size distribution, different grain boundary widths, or different laws of grain growth to be taken into consideration.

2VM2 þ MM1 ¼ 3VM1 þ MM2 þ OO1

Po´´´ 2

J M1 ¼

K 0p1 =W1 y1 Q1 2y1 ð20:81Þ JM ¼ Q1 þ Q2 W1 ð2y1 =W1 Þ þ ðy2 =W2 Þ

JM2 ¼

K 0p2 =W2 y2 Q2 y2 ð20:82Þ JM ¼ Q1 þ Q2 W2 ð2y1 =W1 Þ þ ðy2 =W2 Þ

with JM ¼ JM1 þ JM2 . With this model, it is not possible to express simply the ratio of the thickness of both subscales. However, the ratio of parabolic rate constants can be obtained as Kp0 2 ¼ K 0p1

2 y2 1 þ ðW1 y1 =2y2 W2 Þ y1 1 þ ð2W2 y2 =W1 y1 Þ

ð20:83Þ

When one of the layers is much larger than the other, this expression simplifies to y2 K 0p2 W1 ¼ y1 K 0p1 W2

ð20:84Þ

for example, if y2 y1. It appears that this ratio is directly proportional to the ratio of parabolic rate constants for the

266


growth of each layer alone, that is, roughly proportional to the ratio of the self-diffusion coefficients of the mobile species. This theory has been shown by Garnaud [134] to describe the growth of CuO and Cu2O on Cu, by Garnaud and Rapp [135] to describe the growth of Fe3O4 and FeO on Fe, and by Hsu and Yurek [136] to describe the growth of Co3O4 and CoO on Co.

Generally, experience shows that the oxidation of a metal is often protective during the early stages but the protective properties of the scale can be partially or totally lost during later stages. For example, according to the Haycock–Loriers model [137, 138], the oxidation process involves the concurrent growth of an inner compact layer of MO, controlled by a diffusion mechanism, and its progressive transformation at its outer interface into an external porous layer MO0. The rate of growth of the inner compact layer controlled by diffusional transport is given by ð20:85Þ

where Dm1 is the mass of oxygen in the compact scale, Kp is the parabolic rate constant for growth of the layer, and K1 is the rate constant for its transformation. The growth rate of the outer porous layer controlled by the reaction at its interface with the inner compact layer is given by dðDm2 Þ ¼ fK1 dt

ð20:86Þ

where f is the ratio of the oxygen content in the oxide MO0 to that in the oxide MO. If the scale consists of two layers of the same oxide, f ¼ 1, and the weight gain Dm is given by the rate equation

Kp 2K1 ln 1 ðDm K1 tÞ ¼ 0 Dm þ 2K1 Kp

ð20:87Þ

where the rate constant K1 is characteristic of a transformation in the solid state that does not depend on the oxygen pressure. The function Dm ¼ Dm1 þ Dm2 ¼ F(t) can be approximated by the equation pffiffiffiffiffiffiffi Dm ¼ Kp t þ fK1 t

Δm

Δm1

Kp / 2K1

F2. Paralinear Oxidation

Kp dðDm1 Þ ¼ K1 2 Dm1 dt

Δm

Δm2

Time

FIGURE 20.12. Schematic representation of total weight gain Dm illustrating paralinear rate equation.

This is more likely to be the case if the thermal expansion coefficient of the oxide is much less or greater than unity. The simultaneous oxidation and evaporation leads to the formation of porous and partially porous scales [59]. F3. Stratified Scales Several metals, particularly in columns IV and V (Ti, Zr, Nb, Ta) of the periodic table, form stratified scales during their oxidation at high temperatures as a result of periodic cracking of the growing oxide. From a kinetic point of view, such cracking leads to two types of rate laws (Fig. 20.13): The first one is described by successive parabolic or cubic periods

Type 1 Kinetics transitions

Type 2

ð20:88Þ

which describes paralinear oxidation for which parabolic kinetics predominate during the early stages of oxidation, becoming linear at longer times. This is illustrated in Figure 20.12. Paralinear oxidation is observed during the oxidation of a wide range of metals, especially if mechanical damage occurs to the scale during thermal cycling [139–141].

Time

FIGURE 20.13. Schematic rate laws associated with formation of stratified scales.

OXIDATION RESISTANCE

(type 1), the second one by an initial parabolic or cubic period followed by a near-linear period. Examination of the scale formed on such specimens, either on fractured or polished cross sections, shows that for both type 1 and type 2 rate laws: . .

A compact and adherent scale is formed before the first transition. After the first transition, a porous scale exists which consists of well-defined layers formed essentially parallel to the metal surface; these may be separated by isolated cracks, also parallel to the metal surface, or by connected pores.

Mass transport through the oxide occurs by oxygen, mainly by vacancies V0 (diffusing from the internal to the external interface). Since the scale grows in a confined space at the internal interface, the increase in volume on oxide formation, associated with a Pilling–Bedworth ratio greater than unity, may generate large compressive stresses in the oxide (several gigapascals have been measured). The metal is concurrently submitted to tensile stresses of smaller magnitude (several tens of megapascals) due to its larger thickness. The high compressive stresses in the oxide are probably responsible for the observed tendency for the kinetics to approach a cubic rate law due to an associated reduction in the oxygen diffusion flux. In all cases, oxygen pressure has no influence on the rate constants. As oxide thickness increases, compressive stresses may also increase and result in localized spallation and/or cracking [142, 143]. Then, free access of the gas to a bare metal surface occurs and reoxidaton results. Two cases can be envisaged. In the first, the bare metal is unchanged compared with that at the beginning of oxidation, in terms of composition, microstructure, or mechanical properties. Reoxidation occurs exactly as in the initial period and the rate law is the exact repetition of the first pretransition curve. Kinetics of type 1 (Fig. 20.13) are then observed with a succession of parabolic-type periods. In the second example, the bare metal differs from the initial state, for example, dissolution of oxygen may have occurred, leading to increased hardness and lower creep relaxation rates. The second oxide layer that forms does not then reach the thickness of the first since stress increases at a higher rate and early separation occurs from the metal. In the case of TiO2 growth on titanium at 850 C, for example, the second and all subsequent layers have a thickness of 1–2 mm, whereas the first attains 10–15 mm. The spallation or cracking of these scales does not occur at the same time for all locations on the specimen surface and the resulting law is approximately linear with an increased rate compared to that in the pretransition period. The system can be described as a metal covered by an oxide of statistically constant thickness [8, 144–146].

267

G. OXIDATION RESISTANCE The oxidation of alloys and other metallic materials is a complex process consisting of a large number of phenomena which may, themselves, depend on material composition and environment, that is, temperature, mechanical stress, gas composition, and so on. To obtain effective protection of a metallic material, the oxide formed must lead to a continuous scale (external oxidation) and not to precipitates within the alloy (internal oxidation). Although numerous studies have been devoted to oxidation of alloys, the oxidation mechanisms are often still not fully understood. In fact, the models have usually been developed for the oxidation of the relatively simple example of pure metals and are inadequate to describe alloy oxidation. In order to do this, additional factors have to be taken into account, for example, the different affinity for oxygen of each of the alloy’s constituents, dissolution of oxygen into the alloy, solid solubility between the oxides formed, formation of complex oxides, and different mobilities of the various metal ions in the oxide phases. Moreover, alloy oxidation involves complex processes for which the equilibrium state is reached very slowly or may never be reached. It is unrealistic to try to classify all the various types of alloy oxidation using simple criteria. However, several classification systems have been proposed [67, 147] to account for the diverse morphologies of oxidation scales. They use thermodynamic diagrams and kinetic considerations and, for simplicity, we can consider that three methods are available to protect metals and alloys from high-temperature oxidation: (i) control of the atmosphere, (ii) alloying addition of species more easily oxidized than the base metal (chromium, aluminum) and which form a protective scale during oxidation, and (iii) use of protective coatings deposited by various methods (cementation, plasma spraying, ion bombardment, etc.). Some considerations of these three approaches are presented in the following paragraphs. G1. Atmosphere Control When a process is considered for use in industry, it must be effective, reliable, and economical. For most applications these requirements rule out systems based on high vacuum and the general practice is to use atmospheres derived from fuels. The gases used are therefore mixtures of N2, CO, CO2, H2, H2O, and CH4, which make up the products of combustion of fuels. More recently, atmospheres based on nitrogen have been increasingly used. Starting from the fuel and air, various types of atmospheres can be produced. The main, or common, differentiation is between “exothermic” and “endothermic” atmospheres. The nomenclature is ambiguous and it is well to be clear about its meaning. An exothermic atmosphere is produced exothermically by burning the fuel with measured

268


amounts of air. This type of atmosphere has the highest oxygen potential. An endothermic atmosphere is produced by heating, by external means, a mixture of fuel gas with air over a catalyst to provide a gas containing reducing species. This atmosphere has a low oxygen potential and heat is absorbed during its preparation; hence the atmosphere is described as endothermic. A glance at the analysis of dried, stripped, exothermic atmospheres will confirm that they are predominantly pure nitrogen. Basically, the fuel has been used to remove oxygen from the air. The economics of using nitrogen as a controlled-atmosphere source become more attractive when such factors as safety, reliability, and productivity are considered. Furthermore, the present tendency is to move away from oil toward electricity, in which case nitrogen atmospheres will be particularly attractive. Modern techniques are currently available using carburizing and nitriding systems under vacuum. In these processes of vacuum carburizing and plasma carburizing, the components are heated under vacuum to around 950 C. Methane is leaked into the chamber to a pressure of between 3 and 30 mbar to add carbon to the system. In the absence of a plasma, the methane will only decompose to the extent of about 3%, probably on the surface of the components according to a sequence such as CH4 ! CH3 þ H ! CH2 þ 2H ! CH þ 3H ! C þ 4H ð20:89Þ These reactions may be stimulated to provide 80% decomposition by using a plasma process to excite the methane molecule. In this case, the molecular breakdown may occur in the plasma to produce charged species. Hydrocarbons other than methane may be used as the feedstock. The usual operating sequence involves flushing and evacuation, heating to a temperature under the inert atmosphere, and carburizing for a predetermined time followed by a diffusion anneal in a carbon-free atmosphere. This cycle is designed to provide optimum surface carbon content and carburized depth [9, 10, 12, 13, 148]. The main application of controlled atmospheres is in the area of heat treatment of finished, machined components or of articles of complex shape, which cannot easily be treated subsequently for the removal of surface damage. In this context, the atmosphere is controlled for one of two reasons: to prevent surface reaction or to cause a surface reaction, such as carburizing or nitriding. Prevention or control of oxide layer formation is primarily a matter of controlling the oxygen partial pressure of the atmosphere at a value low enough to prevent oxidation, as described in Section B1. For a metal that undergoes oxidation according to the reaction shown in Eq. (20.7) (Section C), where MO is the lowest oxide of M, the oxygen partial pressure must be controlled so as not to exceed a value pO2 :

pO 2

2 DG ¼ exp RT

ð20:90Þ

where DG is the standard free energy of reaction (20.7). Unfortunately, ðpO2 ÞM-MO is a function of temperature and has lower values at lower temperatures. Thus, if an atmosphere is designed to be effective at high temperatures, it may become oxidizing as the temperature is reduced during cooling. A surface oxide layer may therefore form as the metal is cooled. Although the metallurgical damage to the surface will be negligible, the surface may be discolored, that is, not bright. This condition can be overcome to some extent by rapid cooling or by changing the atmosphere to a lower oxygen partial pressure just before or during cooling. For alloys, the most critical reaction must be considered when deciding on the composition of the atmosphere to be used. For this purpose, the activities of the alloy components must be known since, if the metal M in Eq. (20.7) exists at an activity aM, the corresponding equilibrium oxygen partial pressure will be given by p0O2 : p0O2

1 2 DG ¼ 2 exp RT aM

ð20:91Þ

If the metal activities in the alloy are not known, then, by assuming the solution to be ideal, mole fractions may be used instead of activities to give a value of the oxygen partial pressure at which experiments must be performed to establish the correct atmosphere composition. Low oxygen partial pressures can be provided and, more importantly, controlled by using “redox” gas mixtures. These mixtures consist of an oxidized and a reduced species, which equilibrate with oxygen, for example, CO þ 12O2 ¼ CO2 DG0 ¼ 282200 þ 86:7TJ

ð20:92Þ

from which p00O2 or, more importantly, pCO2 =pCO can be obtained: pCO2 2 2DG0 exp pCO RT

ð20:93Þ

pCO2 DG0 1=2 ¼ pO2 exp pCO RT

ð20:94Þ

p00O2 ¼

Thus, from Eq. (20.93) the ratio of carbon dioxide to carbon monoxide may be calculated for any oxygen partial pressure and temperature. Further discussion of protective atmospheres can be found in references [8], [11], and [36].


G2. Alloying For engineering applications, metals are strengthened and their environmental resistance improved by appropriate alloying. The basic mechanisms operating in pure metal oxidation are also operative in the oxidation of alloys with added complications. These complications include the formation of multiple oxides, mixed oxides, internal oxides, and diffusion interactions within the metals. The effect of alloying on oxidation behavior can be understood first by considering binary alloy model AB consisting of element A, the major component, and B, the minor component [149]. There are two distinct possibilities [150]: (i) One considers element A as more noble and B as more reactive. Thus, at atmospheric pressure of oxygen, B converts to BO, and if the alloy is dilute in B, dispersed precipitates of BO form in A. If the alloy is concentrated in B, a continuous BO scale forms on top of AB. (ii) The other case considers that A and B are both reactive to oxygen with BO more stable than AO. The concentration of B again dictates the oxide morphologies. If the alloy is dilute in B, a stable oxide AO forms as the outer scale. Below the oxide scale, at the AO–alloy interface, the O2 activity is high enough to oxidize B into BO precipitates. If the concentration of B exceeds the critical level required to form a continuous BO scale, then BO forms on top of AB. The oxide growth rate is parabolic, with activation energy characteristic of the growth of BO. The actual rate depends on how protective the BO scale is and on the presence of additional alloying elements. Binary alloys of Ni with Cr or Al form the basis of many oxidation-resistant commercial materials. There are three composition (in wt %) regimes, with distinct oxidation characteristics for each family: Ni–Cr < 10%: Such alloys form an NiO external scale and internal Cr2O3. Ni-30% > Cr > 10%: In such alloys the outer scale consists of NiO on grains and Cr2O3 in grain boundaries. Ni–Cr > 30%: For such alloys the external scale consists of Cr2O3, which is maintained because of the large Cr reservoir. Ni–Al < 6%: Such alloys form an NiO external scale and internal Al2O3 and NiAl2O4. Ni-17% > Al > 6%: In such alloys the outer scale initially consists of Al2O3. However, on continuous exposure, Al depletion occurs in the alloy adjacent to the oxide scale. In the depleted zone, the Al activity is well below the requirement to form continuous Al2O3. Therefore, NiO overtakes Al2O3. The overall result is the formation of a mixture of NiO, NiAl2O4 spinel formed by the combination of NiO and Al2O3. Ni–Al > 17%: For such alloys, the external scale of Al2O3 is maintained because of the large Al reservoir.

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Practical alloys and metallic coatings for high-temperature applications are seldom binary. These alloys are typically Cr2O3 and Al2O3 formers. Silica (SiO2) also forms a protective scale, particularly for refractory metals with which it has better thermal expansion matching. However, SiO2 is not stable at low pressures. It decomposes to gaseous species such as SiO. It also reacts with water vapor at high temperatures, forming Si(OH)4 gas. The use of Cr2O3 scaleforming alloys is limited to temperatures below 1000 C. Volatile CrO3 forms above this temperature in the presence of oxygen due to the reaction 1 2 Cr2 O3

þ 34 O2 ¼ CrO3

ð20:95Þ

Thermodynamic analysis of the foregoing reaction shows 3=4 oxygen partial pressure dependence to be pO2 . Thus, the volatilization becomes important at high oxygen partial pressures. Volatilization of Cr2O3 has been observed even at lower temperatures, between 850 and 900 C. Also, in the presence of water vapor, CrO2(OH)2 forms above the Cr2O3 scale. The scale rapidly vaporizes, resulting in continuing metal recession. The SiO2 scales are also subject to water vapor–enhanced volatility. For high-temperature application, above 1000 C, useful alloys are therefore designed to be Al2O3 formers. In such alloys, the alumina scale can have one of several allotropic forms, depending on the alloy composition and temperature of oxidation. The allotropes include the transient phases g, d, and u which, on thermal exposure, convert into the stable phase a [151]. The composition of the scale formed and its stability depend on the alloy composition, the temperature, and the cyclic nature of thermal exposure [150, 152–156]. Commercial superalloys contain many alloying elements of significance over and above chromium and aluminum. The oxidation behavior of these alloys is very complex and oxidation resistance varies widely, although the general mechanisms described earlier still apply. The complexities arise from significant influence of the individual elemental constituents. Nickel-based superalloys containing elements such as Co, Cr, Al, Ti, W, and Ta exhibit general behavior similar to simple NiCrAl alloys. The cyclic oxidation of superalloys consists of several steps. The process involves an initial transient period during which the oxides of the individual constituents form. These include NiO, CoO, Cr2O3, TiO2, and Ta2O5. Continued oxidation leads to the formation of the most stable oxide, which is Al2O3 for alumina-forming alloys. The parabolic rate constants of the transient oxides are larger than that of the stable oxide. In a cyclic environment the oxide scale cracks and spalls. Aluminum in the alloy diffuses to the oxide–alloy interface to reform the scale. The oxide scale spalling followed by the reformation process continues. An aluminum-depleted zone forms in the alloy below the oxide

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scale. The thickness of the depleted zone depends on the aluminum content of the alloy. Isothermal oxidation of NiCoCrAlY with 12 wt % Al, for example, exhibits depleted zone thickness three times the thickness of the alumina scale [157]. Because of the loss of aluminum, the depleted zone becomes enriched with the other alloying constituents, some of which, having low solubility in the depleted zone, precipitate out in the form of acicular phases. Additionally, nitrogen, which diffuses into the alloy during oxidation, now exceeds its solubility limit and precipitates out as nitrides of such elements as Ti. The new precipitate phases penetrate into the alloy. The alloy is finally depleted of aluminum to below such a critical level that the regeneration of a continuous scale of Al2O3 is no longer feasible. At this stage, for nickel-based alloys, breakaway oxidation starts with the formation of continuous NiO, which is not protective. As a result, oxygen diffuses into the alloy, forming internal oxides of the remaining aluminum, chromium, and other reactive constituents of the alloy. The alloy substrate loses wall thickness and load-bearing capability. To help maintain oxidation resistance, the following steps can be taken: (i) increase aluminum activity as much as possible; (ii) increase aluminum diffusivity to the alloy surface; (iii) inhibit diffusivity of oxygen into the alloy; and (iv) improve the adherence of the alumina scale to the substrate. Similar processes and arguments also hold true for cobaltbased alloys, although some differences exist in the detail. For comparable oxidation resistance, cobalt-based alloys need higher combined aluminum and chromium than do nickel-based alloys. Also, once the protective alumina scale fails, CoO forms on cobalt-based alloys, which spalls off catastrophically as opposed to the gradual failure of NiO formed on nickel-based alloys. The substrate alloy compositions are generally a result of an effort to maximize structural properties such as creep, tensile, and fatigue strength. The alloying constituents, however, affect the oxidation resistance significantly [158–162]. G3. Protective Coatings The focus on the development of substrate alloys discussed in this chapter is generally to achieve high strength, high ductility, and efficient production. Oxidation resistance may not be consistent with achieving these goals. For example, increased Al and Cr result in improved oxidation resistance; however, beyond a certain level, these elements reduce creep strength of the resulting alloys. To achieve both strength and resistance to environmental degradation, the two functions are separated. The load capability is provided by the application of thin coatings with adequate Al and Cr. The thickness of the coating is controlled so that it does not carry any significant load. Depending on the temperature of use, many high-temperature alloys require coatings compatible with its

composition and structural (modulus) and thermal (coefficient of thermal expansion) properties. Diffusion, overlay, and thermal barrier coatings are discussed in the following paragraphs. G3.1. Diffusion Coatings. In diffusion coating, the substrate surface is enriched in an element that will provide high-temperature corrosion resistance. Typical elements are chromium (chromizing), aluminum (aluminizing), or silicon (siliconizing). The substrate is involved in the formation of the coating, and substrate elements are incorporated in the coating; in addition, a diffusion zone is developed in the substrate beneath the coating. Such enrichment not only allows a protective scale to form by selective oxidation but also provides a substantial reservoir of the protective element to delay the inevitable breakaway oxidation to substantially longer times. A wide variety of diffusion coatings are used [163–165]. The most common method of aluminizing is pack cementation, which has been a commercially viable process for many years [166, 167]. Pack cementation and vacuum pack coating techniques can be generalized as methods in which a chemical vapor deposition (CVD) process takes place with the substrate surrounded by a mass of the depositing medium. “Cementation” is a misnomer. The substrate is “packed” in a “cement” consisting of a mixture of the master alloy (the source alloy), a salt as activator, and an inert or reducing hydrogen atmosphere. The coating is carried out over a wide range of pressures from a low, near-vacuum 1–20 torr to nearatmosphere, that is, 360 torr, in the enclosed “retorts” (cf. “reactors” for conventional CVD) [168]. The substrate to be coated is surrounded by the pack; various alignments within this principle are possible depending on the substrate requirements. The first known cementation process was that of Al on steel in 1914 [166, 167]. However, much attention was given to the process, and its variety and development occurred during the mid-1960s to the late 1970s on Ni- and Co-based alloys and iron alloys when the protection of high-temperature gas turbine alloys became paramount together with rocket and space hardware, that is, refractory alloys, mainly Ta, Nb, Mo, and Cr based. The technique itself, in principle, has changed remarkably little since 1914. The composition and quality of the substrate–deposit configuration have been modified as a result of research carried out in the last 10–20 years. Even now, aluminum stands foremost among packcoated deposits, closely followed by Cr, Si, and alloys of Al–Cr, and Al–Si can be coated as one- or two-stage packs [169]. Most of the literature referred to here is on Al, Cr–Al, Ni–Al, and Fe–Al systems. Pack coating is particularly suited to treat large substrates either singly or in bulk and can handle intricate shapes as it is not a


line-of-sight method. Much of the earlier literature is available in references [169]–[171]. Aluminiding was one of the first high-temperature metallic coating systems. Alumina scale-forming coatings are very protective in high-velocity gas turbines. Uncoated alloys containing Al as a minor constituent are rapidly depleted of Al, which diffuses to the surface to maintain alumina formation due to recurring oxide spallation. Diffusion-coated alloys form a surface alloy layer, for example, NiAl and CoAl, which develop the protective alumina layer under more controlled kinetics and reform on spalling without affecting the substrate alloy properties. A renewal is possible before excessive depletion occurs. Brittle intermetallics can be a danger in this system. A diffusion barrier layer between the coating and the substrate is necessary to prevent interdiffusion of elements within the coating system. Multicomponent diffusion coatings (Al, B, Cr, Si, Ti, Zr) to protect machines from liquid Al and for brass corrosion have been used [172]. More studies are needed in this area. On steel, Al coatings are protective up to 500 C, above which brittle intermetallics are formed. Thermal stress-induced cracks then propagate into the substrate metal. Aluminide coatings protect steel from oxidation and corrosion in hydrocarbon and sulfur-containing atmospheres. Aluminized steel is better than stainless steel where oxidation carburization occurs [173]. Chromized steel (diffused in) is resistant to air oxidation up to 700 C. Above 800 C Cr diffuses into the steel, reducing oxidation resistance. At higher temperatures brittle intermetallics form. Good chromized sheet can be bent 180 without damage and is suitable for most firebox and heat exchanger applications up to 600 C. Addition of Al or Si to the chromizing pack process confers oxidation resistance up to 900 C on mild steel, although continuous use at 900 C causes brittle intermetallics and consequent cracking on thermal cycling. On superalloys, alumina-forming diffusion-bonded coatings provide an Al-rich surface to gas turbine environment. Aluminiding superalloys involve formation of more than one phase in the Ni–Al (and Co–Al) systems. Heat treatment is given to stabilize the NiAl phase. NiAl with some Cr and Ti improves the hot corrosion resistance [174]. Oxide particles (e.g., Y2O3) reduce spalling; defects like pinholes, blisters, and cracks may be avoided by a combination of minor additives. Aluminide coatings lack ductility below 750 C and on thermal cycling undergo surface cracking resulting in spalling of the alumina scale. To overcome these two problems, the coating composition was adjusted to embed the brittle b-NiAl or b-CoAl in a ductile g solid-solution matrix. Addition of yttrium improved oxide adherence. Improvements in mechanical properties were achieved by HIP (hot isostatic pressing)–densified, argon-atomized prealloyed powder ingots;

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tensile ductilities of over 20% were produced using the finer precipitate with its better distribution [175]. Studies in which minor amounts of Si, Fe, and Ti were added to improve the scale resulted in little improvement [176]. Platinum electroplate followed by aluminiding gives improved oxidation resistance which offsets the higher cost [177]. Cobalt-based superalloys for higher temperature but less stressed gas turbine vanes have no Al, and this limits the aluminide coating thickness that can be applied without spalling. Superalloy compositions avoiding sigma and other embrittling phases can be destabilized by coating interdiffusion [178]. Structural strengtheners, such as submicrometer oxides in alloys with oxide dispersion strengthening (ODS) and carbides in eutectic alloys with dispersion strengthening (DS), can also limit coating selection [178]. Further research is needed in these areas as higher performance alloys will restrict the choice of acceptable coatings. In a recent five-year period, over 50 production coatings became necessary to replace an original selection of only one to two compositions [179, 180]. G3.2. Overlay Coatings. Overlay coatings are distinguished from diffusion coatings in that the coating material is deposited onto the substrate in ways that give only enough interaction with the substrate to provide bonding of the coating. Since the substrate does not enter substantially into the coating formation, in principle, much greater coating composition flexibility is achievable with overlay coatings. Also, elements, such as Cr, that are difficult to deposit into diffusion coatings can be included in overlay coatings. Although Cr is difficult to incorporate into diffusion aluminide coatings, it is readily incorporated into overlay coatings. Overlay coatings based on the Ni–Cr–Al and Co–Cr–Al systems are commonly used to protect superalloys. Also, small amounts of the reactive elements (e.g., Y, Hf) are routinely incorporated into overlay coatings but are difficult (or sometimes impossible) to incorporate into diffusion coatings. The flexibility in composition of the overlay coatings also allows mechanical properties to be tailored for a given application. Overlay coatings are deposited by physical techniques. The most common are physical vapor deposition (PVD), which includes evaporation, sputtering, and ion plating, and spray techniques (e.g., plasma spraying, flame spraying). Diffusion-type coatings, used successfully on early gas turbines, were tied to the substrate composition, microstructure, and design. Some changes introduced later were (i) superalloy composition, such as reduction in Cr and increase in other refractory metals; (ii) microstructure, in castings with more segregation; and (iii) design, by air cooling and with thin walls (which introduced higher thermal stresses). These changes required coatings that were much more independent of the substrate. Overlay coatings met this necessity.

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Overlay coatings also overcome the process restrictions encountered in diffusion coatings, especially the variants, namely, Cr/Al, Ta þ Cr, or the Pt aluminides, all of which give better stability and oxide hot corrosion resistance than Al alone. MCrAlY compositions (M ¼ Ni, Co, Fe alone or in combination) are the main ones in the series of overlay coatings developed by electron beam evaporated physical vapor deposition (EBPVD) for multiple load. This technique is one of the two most important manufacturing techniques widely used for deposition of thermal barrier coatings (TBCs) onto substrates. Due to its feasibility of varying coating structure, it has great potential for multiple load, for example, to produce functionally graded TBCs. MCrAlYoverlays used in gas turbines are usually Ni and/or Co with high Cr, 5–15% Al, and Yaddition less than about 1% for stability during cyclic oxidation. They are multiphase alloys with ductile matrix (e.g., g Co-Cr) containing a high fraction of brittle phase (e.g., b CoAl). The Cr provides oxidation and hot corrosion resistance, but too much Cr affects substrate-phase stability. The success of most overlay coatings is the presence (and perhaps location) of oxygen-active elements, such as Y and Hf, which promote alumina layer adherence during thermal cycling, giving increased coating protection at lower Al levels. Yttrium appears mostly along grain boundaries if MCrAlY is cast but is homogeneous if plasma sprayed. Thus MCrAlY with 12% Al is more protective than the more brittle diffusion aluminides with 30% Al. Overlay claddings deposited by hot isostatic processing (HIP), electron beam evaporation, or sputtering methods are diffusion bonded at the substrate–coating interface, but the intention here is not to convert the whole coating thickness to NiAl or CoAl. There is, thus, more freedom in coating composition, so that properties can be maximized to the extent required. Compositions based on NiCr, CoCr, NiCrAl, CoCrAl, NiCrAlY, FeCrAlY, and NiCrSi have been used successfully in gas turbine engines. They are generally alumina formers with only 10% Al, unlike the 30% in nickel aluminide coatings. Chromium increases the Al activity allowing this advantage. Higher Al levels cause brittleness and a higher Ductile-Brittle Transition Temperature (DBTT) and the Al levels are generally held below 12% (5–10% preferred). The coatings are also more ductile than NiAl and CoAl and can be rolled and bonded by HIP. In general, NiCrAlY gives best results against high-temperature oxidation whereas CoCrAlY is best for hot corrosion [181]. G3.3. Thermal Barrier Coatings. Thermal barrier coatings are ceramic coatings that are applied to components for the purpose of insulation rather than oxidation protection. The use of an insulating coating coupled with internal air cooling of the component lowers its surface temperature with a corresponding decrease in its creep and oxidation rates. The use of TBCs has resulted in a significant improvement in the efficiency of gas turbines [12].

The earliest TBCs were frit enamels that were applied to aircraft engine components in the 1950s. The first ceramic TBCs were applied by flame spraying and, subsequently, by plasma spraying. The ceramic materials were alumina and zirconia (MgO or CaO stabilized), generally applied directly to the component surface. The effectiveness of these coatings was limited by relatively high thermal conductivity of alumina and problems with destabilization of the zirconia-based materials. Important developments included the introduction of Ni–Cr–Al–Y bond coats and plasma-sprayed Y2O3stabilized zirconia topcoats in the mid-1970s and the development of EBPVD to deposit the topcoat in the early 1980s. Plasma-sprayed TBCs have been used for many years on combustion liners but, with advanced TBCs, vanes, and even the leading edges of blades, can now be coated. The use of TBCs can achieve temperature differentials across the coating of as much as 175 C [182]. Typical systems consist of a nickel-based superalloy substrate coated with M–Cr–Al–Y (M ¼ Ni,Co) or a diffusion aluminide bond coat, which forms an alumina layer [thermally grown oxide (TGO)]. Onto this is deposited a yttria-stabilized zirconia (YSZ) TBC. The TBC can be deposited by air plasma spraying (APS), or EBPVD. The EBPVD coatings are used for the most demanding applications, such as leading edges of airfoils. The APS coating consists of layers of splats with clearly visible porosity and is microcracked. This microcracking is necessary for strain tolerance. The EBPVD coating consists of columnar grains separated by channels similar to the leaders seen in metallic overlay coatings. These channels are responsible for the high strain tolerance of EBPVD TBCs. In summary, a wide variety of coatings are used for oxidation protection at high temperature. It is a huge field dealing with different compositions, production methods, properties, degradation, characterization, repair and function of coatings, and so on. These aspects have been surveyed by Bose [183], Khanna [14], Birks et al. [12], Sato et al. [184], Stern [185], Dahotre and Hampikian [186], Gao [179], and Sch€utze and Quadakkers [180], among others.

H. OXIDATION OF ENGINEERING MATERIALS There is a large spectrum of engineering materials available for applications subject to oxidation in different temperature ranges. This section focuses on laboratory and field experience of many materials that are being considered for applications in various temperature regimes. They range from carbon and Cr–Mo steels to advanced superalloys. At temperatures below 550 C, carbon steel in air shows very little weight gain after exposure for nearly one month. As the temperature is increased to 700 C, the oxidation rate is significantly increased, exhibiting a linear rate of oxidation

OXIDATION OF ENGINEERING MATERIALS

attack. Test results by John [187] showed that carbon steel exhibited about 0.25 mm/year (10 mpy) of oxidation at 604 C; at 650 C, carbon steel exhibited oxidation rates of the order of 1.25 mm/year (50 mpy). The beneficial effects of chromium and silicon additions to carbon steel were reported by Zeuthen [188]. Chromium–molybdenum steels are used at higher temperatures than carbon steel because of higher tensile and creep rupture strengths as well as better microstructural stability. Molybdenum and chromium provide not only solid-solution strengthening but also carbide strengthening. Silicon is very effective in improving the oxidation resistance of Cr–Mo steels, but the most important alloying element for improving oxidation resistance is chromium. The superior oxidation resistance of martensitic and ferritic stainless steels to that of carbon and Cr–Mo steels is well illustrated in the open literature [189]. The 25Cr steel (type 446 ¼ S44600) is the most oxidation resistant among the 400 series stainless steels, due to the development of a continuous Cr2O3 scale on the metal surface. In Fe–Cr alloys, a minimum of approximately 18 wt % Cr is needed to develop a continuous Cr2O3 scale against further oxidation attack [190]. Cyclic oxidation studies conducted by Grodner [191] also revealed that type 446 was the best performer in the 400 series stainless steels, followed by types 430 (S43000; 14–18Cr), 416 (S41600; 12–14Cr), and 410 (S41000; 11.5–13.5Cr). The growth of a thin, adherent (Fe,Cr)2O3 scale as a function of the accumulated isothermal hold time up to 1000 h was observed by Walter et al. [192] at 650 C in air during cycling from 650 to 300 C. The 300 series austenitic stainless steels have been widely used for high-temperature components because they exhibit higher elevated-temperature strength than do ferritic stainless steels. Furthermore, they do not suffer 475 C embrittlement or ductility loss problems in thick sections and in heat-affected zones as do ferritic stainless steels. Nevertheless, some austenitic stainless steels can suffer some ductility loss upon longterm exposure to intermediate temperatures (e.g., 540–800 C) due to sigma-phase formation [193]. Moccari and Ali [194] observed the beneficial effects of nickel in improving the oxidation resistance of austenitic stainless steels. In evaluating materials for automobile emission control devices, such as thermal reactors and catalytic converters, Kado et al. [195] carried out cyclic oxidation tests on various stainless steels. In cyclic oxidation tests performed in still air at 1000 C for 400 cycles (30 min in the furnace and 30 min out of the furnace), types 409 (S40900; 12Cr), 420 (S42000; 13Cr), and 304 (S30400; 18Cr–8Ni) suffered severe attack. Type 420 (S42000; 13Cr) was completely oxidized after only 100 cycles, although the sample did not show any weight changes. Alloys that performed well under these conditions were types 405 (S40500; 14Cr), 430 (S43000; 17Cr), 446 (S44600; 25Cr), 310 (S31000; 25Cr–20Ni), and DIN 4828 (19Cr–12Ni–2Si).

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When cycled to 1200 C for 400 cycles (30 min in the furnace and 30 min out of the furnace), all alloys tested except F-1 alloy (Fe–15Cr–4Al) suffered severe oxidation attack. This illustrates the superior oxidation resistance of alumina formers (i.e., alloys that form Al2O3 scales when oxidized at elevated temperatures). Their data also illustrate that, for temperatures as high as 1200 C, Cr2O3 oxide scales can no longer provide adequate oxidation resistance. Oxidation data generated in combustion atmospheres are relatively limited. No systematic studies have been reported that varied combustion conditions, such as air-to-fuel ratios. In combustion atmospheres, the oxidation of metals or alloys is not controlled by oxygen only. Other combustion products, such as H2O, CO, CO2, N2, hydrocarbon, and others, are expected to influence oxidation behavior. When air is used for combustion, nitride formation in conjunction with oxidation can occur in combustion atmospheres under certain conditions. The presence of water vapor can also be an important factor in affecting oxidation behavior of alloys, as discussed later in this section. Manufacturing processes can greatly influence the surface chemistry of an alloy product. Stainless steels can be finished into the final product by bright annealing (i.e., annealing is performed in a protective atmosphere, such as hydrogen environment or dissociated ammonia environment). This process generally produces a product with minimal depletion of chromium at or near the surface. On the other hand, when the alloy product is finished by black annealing (i.e., annealing is performed in air or combustion atmosphere in the furnace) and followed by acid pickling, there is a good chance that the alloy product may exhibit surface depletion of chromium. This is particularly important for thin-gage sheet products or thin tubular products [196]. Some stainless steel producers may manufacture stainless steels at the bottom of the specification range for key alloying elements, such as chromium, to reduce materials cost. Accordingly, the chromium content can be insufficient to maintain a continuous chromium oxide scale during prolonged service or when subjected to thermal cycling or overheating conditions, thus promoting breakaway oxidation. The oxidation resistance of these “lean” stainless steels can be further degraded by the surface depletion of chromium resulting from manufacturing processes that may involve excessive pickling after “black” annealing (annealing in air or combustion atmosphere), during successive reductions in cold rolling in flat product manufacturing, or pilgering in tubular manufacturing. The chromium concentration at the surface of such a product, particularly a thin-gage sheet or tube, may be too low to form or maintain a continuous chromium oxide scale during service. As a result, iron oxides and isolated nonprotective Fe–Cr oxide nodules can develop on the metal surface, thus resulting in breakaway oxidation, as discussed later.

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Some commercial electrical resistance heating elements are made of Fe–Cr–Al alloys, such as Kanthal alloys, which rely on the formation of the Al2O3 scale for applications up to 1400 C [197]. For example, some of the Kanthal alloys that are available in wire, strip, and ribbon product forms are Kanthal A-I (K92500; Fe–22Cr–5.8Al), AF (Fe–22Cr– 5.3Al), and D (Fe–22Cr–4.8Al). Since these wrought alloy products are essentially ferritic alloys, they exhibit low creep rupture strengths when the temperature exceeds 650 C and cannot be used for high-temperature structural components. Thus, the electrical resistance heating elements made of these alloys must be properly supported to avoid creep deformation, such as sagging. These Kanthal wires can be used in arc or flame spraying to produce an oxidation-resistant coating or in weld overlay cladding by using a gas metal arc welding (GMAW) process. A powder metallurgy (P/M) process was used to produce a new alloy product, Kanthal APM, reported to exhibit improved creep rupture strength [198]. Other commercial Fe–Cr–Al alloys include ALFA-ITM (Fe–13Cr–3Al), ALFA-IITM (Fe–13Cr–4Al), and ALFA-IVTM (Fe–20Cr–5Al–Ce) developed by Allegheny Ludlum [199] and Fecralloy (Fe–16Cr–4Al–0.3Y) developed by Atomic Energy Authority [200]. As the nickel content in the Fe–Ni–Cr system increases from austenitic stainless steels to a group of iron-based alloys with 20–25Cr and 30–40Ni, the alloys become more stable in terms of metallurgical structure and more resistant to creep deformation (i.e., higher creep rupture strengths). In general, this group of alloys also exhibits better oxidation resistance. Some of the wrought alloys in this group are 800H/800HT (Fe–21Cr–32Ni–Al–Ti), RA330 (N08330; Fe–19Cr–35Ni–1.2Si), HR120 (Fe–25Cr–37Ni–0.7Nb-N), AC66 (N33228; Fe–27Cr–32Ni–0.8Nb–Ce), 353MA (Fe–25Cr– 35Ni–1.5Si–Ce), and 803 (Fe–26Cr–35Ni–Al–Ti) [201]. In many Ni–Cr alloys, many alloying elements, such as those for solid-solution strengthening (e.g., Mo, W) and precipitation strengthening (e.g., Al, Ti, Nb), are added to the alloys to provide strengthening of the alloy at elevated temperatures. Many of these alloys are commonly referred to as “superalloys.” The superalloys also include oxide dispersion–strengthened (ODS) alloys, which are briefly discussed later. Similar to Fe–Cr–Al alloys, aluminum is also used as an alloying element in Ni–Cr alloys to improve the oxidation resistance. Although a Ni–Cr alloy generally requires a minimum of 4% Al to form a protective Al2O3 scale, the addition of less than 4% Al can significantly improve the oxidation resistance of the alloy. Alloy 601 (N06601), with only about 1.3% Al, shows excellent oxidation resistance [202]. Nickel–chromium alloys containing about 4% Al or higher form a very protective Al2O3 scale when heated to very high temperatures; for example, Lai [203] compared alloy 214 (Ni–16Cr–4.5Al–Y) with alloy 601 and alloy 800H

in cyclic oxidation tests performed in still air at 1150 C with specimens cycling to room temperature once a day except weekends. Alloy 214 showed essentially no weight loss after 42 days of testing, whereas alloy 601 showed a linear weight loss. For applications at high temperatures, many superalloys contain numerous alloying elements for increasing the elevated-temperature strength of the alloy. Molybdenum and tungsten are common alloying elements for providing solidsolution strengthening for increasing the creep rupture strength of the alloy. Two iron-based superalloys, Multimet alloy (R30155; Fe–20Ni–20Co–21Cr–3Mo–2.5W–1.0Nb þ Ta) and alloy 556 (R30556; Fe–20Ni–18Co–22Cr– 3Mo–2.5W–0.6Ta–0.02La–0.02Zr), are good examples. However, the oxides of both molybdenum and tungsten (MoO3 and WO3) exhibit high vapor pressures at very high temperatures. Multimet alloy showed rapid oxidation attack at 1150 and 1200 C, with specimens completely consumed at both temperatures. However, formation of the volatile oxides of MoO3 and WO3 can be minimized by modification of some key alloying elements in Multimet alloy. The development of alloy 556 was aimed at improving the oxidation resistance of Multimet alloy without losing the elevatedtemperature strength by modifying the Multimet alloy composition. The modification involved a slight increase in chromium, a decrease in cobalt, replacement of niobium with tantalum, and addition of a rare earth element, lanthanum, and a reactive element, zirconium, but the concentrations of molybdenum and tungsten were not changed. The result was a much more oxidation-resistant alloy, alloy 556, at 1095 and 1150 C, although rapid oxidation nevertheless occurred at 1200 C. Cobalt-based alloys with tungsten, such as alloy 188 (R30188; Co–22Cr–22Ni–14W–0.04La), alloy 25 (R30605; Co–20Cr–10Ni–15W), and alloy 6B (Co–30Cr– 4.5W–1.2C), showed rapid oxidation at 1205 C. A cobaltbased alloy, alloy 150 (Co–27Cr–18Fe), containing no tungsten also showed rapid oxidation attack at 1205 C. Again, the oxidation of a cobalt-based alloy can be significantly improved with some modification of alloying elements. Alloy 25 with 15% W exhibits excellent creep rupture strength at high temperatures. However, because of the high level of tungsten, the alloy shows high oxidation rates at very high temperatures, such as 1095 and 1150 C. With a slight increase in chromium and nickel along with the addition of lanthanum, the result of the modification was alloy 188, which has significantly better oxidation resistance than alloy 25 at 1095 and 1150 C. Oxide-dispersion-strengthened alloys use very fine oxide particles that are uniformly distributed throughout the matrix to provide excessive strengthening at very high temperatures. These oxide particles, typically yttrium oxide, do not react with the alloy matrix, and so no coarsening or dissolution occurs during exposure to very high temperatures, thus

OXIDATION OF ENGINEERING MATERIALS

maintaining the strength of the alloy. This group of superalloys is produced using specialty powders that are manufactured by the mechanical alloying process. These powders are essentially composite powders with each particle containing a uniform distribution of submicrometer oxide particles in an alloy matrix. The process of producing these ODS powders involves repeated fracturing and rewelding of a mixture of powder particles in vertical attritors or horizontal ball mills [204]. A vertical attritor is a high-energy ball mill (like the conventional horizontal ball mill) in which the balls and the metal powders are charged into a stationary vertical tank and are agitated by impellers rotating from a central rotating shaft. The shaft turns, but the jar stays put. As the balls are stirred, they fall on the contents and grind whatever is between the balls. Alloy powders are then canned, degassed, and hot extruded followed by hot working and annealing to produce a textured microstructure. Alloys are available in mill products such as bar, plate, and sheet or custom forgings. Oxidation of alloys can significantly increase under highvelocity gas streams, as in, for example, combustors and transition ducts in gas turbines. These components are also subject to severe thermal cycling, particularly gas turbines in airplane engines. Laboratory burner rigs have been developed to evaluate the type of oxidation, often referred to as “dynamic oxidation,” under conditions of very high gas velocities. Some of these dynamic oxidation burner rigs are described elsewhere [204–206]. Hicks [207] performed dynamic oxidation tests with 170 m/s gas velocity at 1100 C with 30-min cycles for several wrought chromia former superalloys and an ODS alumina former (MA956). Alumina former MA956 was found to be considerably better than chromium formers, such as alloys 191, 86, 617, 188, and 263. MA956, along with some ODS alloys, was studied by Lowell et al. [204] with 0.3 Mach gas velocity at 1100 C with 60-min cycles. ODS alloys included in the study were MA956 (Fe–19Cr–4.4Al–0.6Y2O3), HDA8077 (Ni–16Cr– 4.2Al–1.6Y2O3), TD-NiCr (Ni–20Cr–2.2ThO2), and STCA264 (Ni–16Cr–4.5Al–1Co–1.5Y2O3). Also included in the study was a PVD coating of Ni–15Cr–17Al–0.2Y on MAR-M-200 alloy (Ni–9Cr–10Co–12W–1Nb–5Al–2Ti). MA956 and HDA8077 as well as PVD Ni–Cr–Al–Y coating were found to perform well. No explanation was offered for STCA264, which did not perform as well as HDA8077, although both alloys had similar chemical compositions. In Fe–Cr, Fe–Ni–Cr, Ni–Cr, and Co–Cr alloy systems, the formation of an external Cr2O3 oxide scale provides the oxidation resistance for the alloy. The growth of the Cr2O3 oxide scale follows a parabolic rate law as the exposure time increases. As the temperature increases, the oxide scale growth rate also increases. The growth of the Cr2O3 scale requires that a continuous supply of chromium from the alloy interior diffuses to the oxide–metal interface. Continued

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oxidation can eventually deplete chromium in the alloy matrix immediately under the oxide scale. When the chromium concentration in the alloy matrix immediately beneath the oxide scale is reduced to below a critical concentration, the alloy matrix no longer has adequate chromium to reform a protective Cr2O3 oxide scale when the scale cracks or spalls due to oxide growth stresses or thermal cycling. Once this occurs, fast-growing, nonprotective iron oxides, or nickel oxides, or cobalt oxides (i.e., oxides of base metal) form and grow on the alloy surface. Breakaway oxidation initiates, and the alloy begins to undergo oxidation at a rapid rate. The alloy thus requires the level of chromium immediately under the chromium oxide scale to reheal [206–211]. To prolong the time for initiation of breakaway oxidation, it is necessary to have an adequate reservoir of chromium immediately below the oxide scale to provide adequate chromium to maintain a protective chromium oxide scale or to reheal the oxide scale after local cracking or failure. For alumina formers such as Fe–Cr–Al alloys and Fe– Cr–Al- and Ni–Cr–Al-based ODS alloys, breakaway oxidation occurs when aluminum concentration under the Al2O3 scale has been reduced to a critical level such that healing of the Al2O3 is no longer possible, thus resulting in the formation of nonprotective, fast-growing oxides of base metals (e.g., iron oxides or nickel oxides). The breakaway oxidation due to rapid growth of iron oxides or nickel oxides becomes essentially a life-limiting factor. This critical aluminum concentration was found to be about 1.0–1.3% for Fe– Cr–Al-based ODS alloys (e.g., MA956, ODM751) at 1100–1200 C [210, 211]. These values were obtained from foil specimens (0.2–2 mm thick) tested in still air at 1100–1200 C. For the non-ODS Fe–20Cr–5Al alloy, this critical aluminum concentration was found to be higher (about 2.5%) at 1200 C [210]. Since breakaway oxidation is related to the aluminum reservoir in the alloy, this reservoir is a critical issue when the component is made of thin sheet or foil. Because of excellent oxidation resistance at very high temperatures, there is increasing interest in considering alumina formers for products that require thin foils, such as honeycomb seals in gas turbines, metallic substrates for automobile catalyst converters, and recuperators in microturbines. For alumina formers to improve their resistance to breakaway oxidation, yttrium is frequently used to increase the adhesion of the aluminum oxide scale. Other alloying elements that are known to increase the adhesion of the aluminum oxide scale include zirconium and hafnium. Quadakkers [212] shows that both MA956 (Fe–20Cr–4.5Al– 0.5Y2O3) and Aluchrom (Fe–20Cr–5Al–0.01Y) exhibited much more cyclic oxidation resistance than Fe–20Cr–5Al when tested at 1100 C in synthetic air with an hourly cycle to room temperature. Addition of Y2O3 to an alumina former has a similar beneficial effect as yttrium added as an alloying element.

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Klower and Li [213] studied the oxidation resistance of Fe–20Cr–5Al alloys in 10 different compositions containing various amounts of yttrium ranging from 0.045 to 0.28%. All 10 compositions contained 0.002% S, and 8 compositions contained 0.04–0.06% Zr with two compositions containing no zirconium. The cyclic oxidation tests were performed at 1100 and 1200 C, respectively, with each cycle consisting of 96 h at temperature and rapid air cooling to room temperature. These authors concluded that the yttrium addition of about 0.045% was sufficient to prevent the oxide scales from spalling, and when the yttrium concentration was increased to more than 0.08%, substantial internal oxidation could occur, resulting in rapid metal wastage [213]. Sulfur in the alloy is known to play a very significant role in adhesion of the aluminum oxide scale to the alloy substrate for alumina formers. The role of yttrium is believed to prevent the preferential segregation of sulfur in the alloy to the scale–metal interface to weaken the adhesion of the oxide scale [214–216]. Reducing the concentration of sulfur in a Ni–Cr–Al alloy can significantly improve the oxidation resistance of the alloy. Smeggil [217] compared cyclic oxidation resistance between the normal purity Ni–Cr–Al alloys (approximately 30–40 ppm S) with the high-purity Ni–Cr–Al alloys (approximately 1–2 ppm S), showing a significant improvement in cyclic oxidation resistance when sulfur in the alloy was significantly reduced. Also demonstrated was the beneficial effect of yttrium addition to the normal purity Ni–20Cr–12Al alloy, showing significant improvement in the cyclic oxidation resistance of the alloy without reducing the sulfur content in the alloy. Sulfur was found to segregate to the oxide–alloy interface during oxidation in Fe–Cr–Al alloys [218, 219]. Yttrium is believed to tie up sulfur at the oxide–metal interface, thus improving the oxide scale adhesion [217]. There are some industrial applications that require thingage sheet materials or thin foils for construction of critical components. As the component thickness decreases, oxidation becomes a major factor that limits service life. When the component is made of thin foil, prolonging the incubation time before initiation of breakaway oxidation is the controlling factor for extending the service life of the component. Thus, as applications are being pushed toward higher and higher temperatures, alloys that form aluminum oxide scales can offer significant advantages in performance over those alloys that form chromium oxide scales [220–222]. In high-temperature combustion atmospheres, water vapor is invariably present in the environment. The effect of water vapor on the oxidation of alloys is an important factor in the alloy selection process. Most oxidation data are generated in laboratory air, which generally contains low levels of water vapor [223–225]. The effect of water on oxidation, and its detrimental effect, has been established by Onal et al. [226] and many others [2, 4, 12, 16, 170, 183, 227].

As temperature increases, metals and alloys generally suffer increasingly higher rates of oxidation. When the temperature is excessively high, metals and alloys can suffer rapid oxidation. There is, however, another mode of rapid oxidation that takes place at relatively low temperatures, often referred to as “catastrophic oxidation,” associated with the formation of a liquid oxide that disrupts and dissolves the protective oxide scale, causing the alloy to suffer rapid oxidation at relatively low temperatures. This phenomenon has been observed by, for example, Meijering [228], Brennor [229], Sawyer [230], Brasunas [231], and Sequeira [232]. The most effective way to alleviate the potential catastrophic oxidation problem is to avoid a stagnant condition of the gaseous atmosphere.

I. CONCLUSIONS This chapter is a summary of the main factors for determining the nature and extent of gas–metal reactions, the comprehension of which is of paramount importance to understand the subject. Some of these factors are metal structure, oxide structure, metal diffusion, oxygen or metal diffusion at grain boundaries or on the surface, metal and oxide volatility, oxygen solution in metal, cracking, spalling, blistering, sintering, oxide nucleation, electronic conductance, oxide adhesion, and oxide plasticity.

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21 THERMOCHEMICAL EVALUATION OF CORROSION PRODUCT STABILITIES FOR ALLOYS IN GASES AT HIGH TEMPERATURE W. T. THOMPSON Center for Research in Computational Thermochemistry, Royal Military College of Canada, Kingston, Ontario, Canada

R. C. JOHN Shell International E&P, Inc. Houston, Texas

A. L. YOUNG Humberside Solutions Ltd., Toronto, Ontario, Canada

A. Introduction B. Thermodynamics B1. Gas phase B2. Corrosion product formation C. Kinetics D. Software E. Conclusion Acknowledgment References

kinetics associated with growth can be treated. Thermochemical analysis of the potential corrosion products can help identify the corrosion mechanisms important for the process conditions and also suggest which alloy types might have sufficient corrosion resistance based upon knowledge of how corrosion products influence corrosion rates. Examples of processes that corrode metals in high-temperature gases include coal gasification, crude oil distilling, steam-methane re-forming, catalytic cracking, fluid bed coal combustion, petroleum coking, combustion, incineration, hydrocracking, naphtha cracking to ethylene, and other energy conversion processes. All involve the presence of gases containing H–O–S–Cattemperaturesof600–1300 K(300–1000 C)[1].

A. INTRODUCTION B. THERMODYNAMICS The understanding of corrosion of metals and alloys in hot gases can be assisted by analysis of the thermodynamics of the formation of corrosion product phases involving elements in the alloy and species in the gas. Since the service temperatures and pressures in high-temperature applications may vary over large ranges and the gas phase may involve many species, the stable corrosion products are only known with certainty for applications of specific alloys. It is often helpful to employ a multielement stability (or phase) diagram, which considers the thermodynamic properties of the possible compounds, to establish the most stable corrosion products before the

B1. Gas Phase In many processes, the nominal gas composition may not represent the composition present at the corroding surface. In the process of sampling gases containing H, O, S, and C at high temperature and pressure, part of the water vapor may condense during subsequent cooling and pressure reduction. Furthermore, in mixtures rich in CO, there is the possibility of carbon precipitation: 2CO ¼ CO2 þ C

ð21:1Þ


281

282

THERMOCHEMICAL EVALUATION OF CORROSION PRODUCT STABILITIES FOR ALLOYS IN GASES AT HIGH TEMPERATURE

resultant from a shift in the equilibrium with falling temperature given approximately by log

" # ðPCO2 ÞðaC Þ 2

ðPCO Þ

¼

8916 9:15 T

ð21:2Þ

Even if the concentrations of the most abundant gases change little with temperature, the partial pressures of the constituent elements such as O2 or S2 may vary considerably. It may be necessary to infer, from a gas composition at ambient temperature and pressure, what the gas composition is likely to be at the service condition for the alloy. Table 21.1 shows the computed equilibrium composition at 900 K (627 C) and 1100 K (827 C) for a gas with an initial composition of 35% CO, 25% H2, 20% CO2, 19% H2O, and 1% H2S. In both cases, the total pressure is 30 atm. This type of gas composition might be found in coke/coal/oil gasification, Flexicoking (trade name of Exxon Corporation), catalytic cracking in a petroleum refinery, steam-methane re-forming, or instances of fuel-rich conditions of fossil fuel combustion. The nominal (molar) composition may be regarded as the reactants in a chemical process leading to the equilibrium products. In the product equilibrium, the mole fraction is given adjacent to each species considered. The total number of moles of gas appears near the bottom of each listing and is, of course, larger for the higher temperature reflecting the greater degree of dissociation. Although graphite is a potential phase in both figures, it only appears at the lower temperature; at the higher temperature the activity of carbon with respect to graphite is less than 1. The equilibrium was established by the Gibbs energy minimization [2, 3]. In this versatile computational procedure, the molar concentration of candidate equilibrium species (containing some or all of the elements in the nominal gas) are systematically varied subject to the condition of mass balance for each element. For each such mixture, the total (relative) Gibbs energy, Gt, is calculated by evaluating the following function: Gt ¼

gases X i¼1

X solids ni fDGi þ RT ln½ðXi ÞðPÞg þ ni DGi i¼1

ð21:3Þ

where DGi are the standard Gibbs energies of formation per mole of each species in the equilibrium gas mixture, R is the gas constant, ni is the number of moles, Xi is the mole fraction of species in the gas phase at absolute temperature, T, and total pressure, P. The equilibrium is found when successive (and progressively smaller) adjustments in molecular concentrations provide negligible reduction in Gt. The equilibrium partial pressure of the elemental species (PH2 ; PO2 ; PS2 ) and the activity of carbon (graphite, aC) are a by-product of the computation.

B2. Corrosion Product Formation Having established the equilibrium gas composition at the temperature of interest, it is now possible to consider the reaction of that gas with the elements in the alloy. Consider the case of Inconel 671 (trade name of Huntington Alloys International), which is nearly a binary alloy containing 54% Ni and 46% Cr. This alloy is not widely used but is chosen for this example to illustrate the concepts used to calculate corrosion product stabilities. The most stable chemical phase of nickel can be determined by computing the Gibbs energy change for all possible reactions of nickel with hydrogen, oxygen, sulfur, and carbon, assuming that the composition of the voluminous gas phase is not influenced by reactions with a relatively minor mass of the alloy. The reactions, all based on 1 mol of nickel reacting to form a series of Ni compounds with the general formula NiHaObScCd, have the form: Ni þ

a 2

H2 þ

c b O2 þ S2 þ ðdÞC ! ðNiHa Ob Sc Cd Þ 2 2 ð21:4Þ

where the constants a, b, c, and d are 0. When the Gibbs energy change is computed, allowance is made for PH2 ; PO2 ; and PS2 , and the activity of C appropriate to the gas mixture at the temperature and total pressure involved using the equation a b RT ln PO2 RT ln PH2 2 2 c RT ln PS2 ðdÞRT ln aC 2

DG ¼ DG

ð21:5Þ

where DG is the standard Gibbs energy of formation for the nickel-containing compounds (per mole of Ni) obtained from tables or databases of thermochemical properties. The reaction with the most negative Gibbs energy change identifies the most stable compound containing nickel coexisting in direct contact with the gas phase. In order to evaluate the potential of changes in the gas composition to alter the most stable nickel compound, a diagram may be constructed by systematically varying the partial pressures of two particular elements while holding the partial pressure or activity of the third and fourth elements constant. Such a predominance or phase stability diagram is shown in Figure 21.1 for 1100 K (827 C) [2, 4]. Note that the suggested methodology for the diagram construction does not begin with potentially false suppositions of coexistence of particular combinations of nickel containing phases as a basis for developing equations for each phase boundary. Therein lies the power of the suggested computational procedure. Precision in locating the boundary for graphical purposes simply involves selecting suitably small steps in the variation of the partial pressures for the species used for the

THERMODYNAMICS

FIGURE 21.1. The Ni-H, O, S, C predominance diagram at 1100 K (827 C) prepared using the FACT system [2]. The coordinate is consistent with Table 21.1.

axes. Many computations of Gibbs energy can be eliminated following the realization that phase fields must be contiguous. The computational power of a personal computer is more than ample to develop diagrams as in Figure 21.1 in a few seconds [2]. When it is more informative to employ the dominant gas species for the axes and fixed partial pressures in the diagram, a variation on the foregoing methodology is invoked Suppose, for example, log PCO2 is used as oidinate and log PCO as abscissa. The equilibrium constant for the process CO þ 12O2 ¼ CO2 PCO2 14; 764 ¼ log 4:537 ðPCO ÞðPO2 Þ T

ð21:6Þ

ð21:7Þ

enables PO2 to be found for any pair of values of PCO and PCO2 . Similarly, the carbon (graphite) activity can be found using Eqs. 21.1 and 21.2. If, for the diagram development, the PH2 O and PH2 S were fixed, then PH2 could be found from the equilibrium constant for H2 þ 12O2 ¼ H2 O

ð21:8Þ

and PS2 from the equilibrium constant for H2 þ 12S2 ¼ H2 S

ð21:9Þ

283

FIGURE 21.2. The Ni-H, O, S, C predominance diagram at 1100 K (827 C). The diagram is similar to Figure 21.1 but a different selection of axes and constant conditions has been made. The coordinate is consistent with Figure 21.1 and Table 21.1.

These preliminary steps reduce the computations used to construct Figure 21.2 to those of Figure 21.1. The two diagrams are consistent in indicating that Ni3S2 (liquid) is the most stable phase of nickel in equilibrium with the gas. Comparable points, representative of the equilibrium gas in Table 21.1, are shown on both figures. The two diagrams are not topologically the same with respect to triple points and adjacent phase fields since the chemical potentials of all elements in the gas phase are only the same at the coordinates marked. In the case of corrosion of Inconel 671, it is of course necessary to also consider the chemical compounds of Cr. By procedures similar to that described for Ni, a diagram similar to Figure 21.2 may be developed for Cr. This diagram superimposed on that for Ni is shown in Figure 21.3. The overlapping stability fields give rise to the double labeling of each area on the resultant diagram. For the conditions of temperature, total pressure, and gas compositions covered by Figure 21.3, there are no stable compounds containing both Ni and Cr and little or no mutual solubility of the various nickel- or chromium-containing compounds. Thus the superimposition of Ni and Cr stability diagrams leads to the conclusion, for the gas with the equilibrium composition in Table 21.1 at 1100 K, that liquid Ni3S2 and Cr2O3 are the most stable corrosion products. As evidence that Ni3S2 and Cr2O3 cannot react to form, for example, NiCr2O4 in contact with the equilibrium gas

284


composition in Table 21.1 at 1100 K, it is necessary to evaluate the Gibbs energy change for the reaction of Ni3S2 and Cr2O3 to form NiCr2O4 (balanced using CO2, CO, H2O, and H2S). Ni3 S2 þ 3Cr2 O3 þ CO2 þ 2H2 O ! 3NiCr2 O4 þ 2H2 S þ CO ð21:10Þ

FIGURE 21.3. The Ni, Cr-H, O, S, C predominance diagram at 1100 K (827 C). The diagram in this case is the superimposition of Figure 21.2 and a similar one for Cr since, within the field of view, there are no phases that contain both Ni and Cr (such as NiCr2O4). The placement of the coordinate consistent with Table 21.1 shows that liquid Ni3S2 is a stable phase.

In finding the Gibbs energy change at 1100 K [2], allowance is made for the partial pressures of the gas species at the composition of interest using an equation similar to Eq. (21.5). The positive value of 146 kJ/mol affirms the inability of the sulfide and oxide to react; that is, the stability of Ni3S2 and Cr2O3 in direct contact in this particular atmosphere. Although the superimposition of Ni and Cr stability diagrams provides (by good fortune, in this particular case) a suitable corrosion product stability diagram, superimposition is not sufficiently general to create a reliable diagram in all cases. A more universally applicable approach is a necessity. To take the case of Inconel 671, the construction of a composite stability diagram for two (or more) alloying elements can be undertaken by finding the Gibbs energy changes for all reactions of the type: a b c H2 þ O2 þ S2 þ ðdÞC ð1 rÞNi þ r Cr þ 2 2 2 ! m½Niw Crx ðH; O; S; C; Þ þ n½Niy Crz ðH; O; S; C; Þ ð21:11Þ

TABLE 21.1. Equilibrium Computations for the Initial Gas Composition of 35 Mol CO, 25 Mol H2, 20 Mol CO2, 19 Mol H2O, and 1 Mol H2S, Shown as Mole Fractions Adjacent to the Given Species Species CO CO2 H2 H2O CH4 H2S COS H2CO C2H4 C2H6 H2S2 CH3SH S2 SO2 SO3 O O2 Total moles of gas Activity of C Moles of C

(900 K, 30 atm) Xi

(1100 K, 30 atm) Xi

0.050 0.431 0.091 0.335 0.079 0.013 0.230 10 3 0.225 10 6 0.897 10 7 0.189 10 5 0.323 10 6 0.291 10 6 0.323 10 8 0.217 10 9 0.171 10 19 0.209 10 23 0.457 10 23 75.3 1 12.8

0.286 0.267 0.206 0.192 0.039 0.010 0.455 10 3 0.214 10 5 0.821 10 6 0.612 10 6 0.420 10 6 0.228 10 6 0.311 10 7 0.110 10 8 0.841 10 18 0.101 10 18 0.493 10 19 92.8 0.838 0

In this generalized methodology [5], following the constraints of the phase rule, groups of compounds equal in number to the components in the alloy phase are considered. All the candidate corrosion product compounds in the group must contain some or all of the alloying elements and the molar amounts given by m and n must both of course be nonnegative for the proportion of elements in the alloy given by r. With all equations balanced in terms of 1 mol of alloy with appropriate zero or positive coefficients a, b, c, and d, the Gibbs energy change for each is found in a manner similar to that for reactions of the type given by Eq. (21.4). The corrosion product pair that is most stable is associated with the reaction with the most negative Gibbs energy change. The topology of such diagrams is sensitive to the value of r. Under more oxidizing conditions than those usually associated with coal gasification applications, NiCr2O4 may be a stable corrosion product. Clearly, whether NiCr2O4 coexists with a Cr- or a Ni-containing compound depends on whether r is greater or less than 23. C. KINETICS The qualitative characterization of corrosion rate in hightemperature systems may sometimes be gleaned from a

SOFTWARE

combination of stability diagrams and corrosion rate data. Clearly, the stability diagram for Inconel 671 at 1100 K shown in Figure 21.3 is revealing in terms of indicating liquid sulfide corrosion products, which invariably have high corrosion rates, in contact with the equilibrium gas compositions represented in Table 21.1. Liquid Ni3S2, as a stable corrosion product (coexisting with solid Cr2O3), would be expected to corrode rapidly in the gas at 1100 K. At the lower temperature (900 K) associated with Figure 21.4, where Ni3S2 is solid, somewhat slower corrosion rates might be expected. Thermochemical information is valuable in organizing and interpreting data on corrosion at high temperatures, for design and selection purposes, in combination with corrosion rates measured in actual exposure testing. To quantitatively represent isothermal corrosion testing, it is possible to provide empirical constants in an equation of the form p ¼ ktn

ð21:12Þ

where p is the penetration (or depth of loss of internally oxidation free alloy) and k and n are characteristic parameters established by a statistical treatment of sufficient test data [6]. For a pure metal with only one corrosion product, a value of n near 1 may indicate that the superficial scale is not protective but rather spalls at a constant rate perhaps due to

285

excessive volume in comparison to the volume of metal replaced (Pilling–Bedworth ratio [7]). In other cases, corrosion rate is limited by arrival of a corrosion species or removal of a volatile corrosion product compound. Protective corrosion products must have a value of n < 1. If the growth of the scale is characterized by a constant diffusion coefficient in the alloy or corrosion product compound, n should be 12. This is sometimes termed parabolic growth. Defects in the crystallinity of the scale, particularly as affected by impurities or alloying elements incorporated into the growing scale, may lead to values other than 12. Many corrosion–resistant alloys often have both a value of n of 12 and a small value for k. The parameter k is temperature dependent and is sensitive to the equilibrium partial pressure of the elements in the gas phase as well. When Inconel 671 is exposed to oxidizing and sulfidizing atmospheres such that Ni3S2 and Cr2O3 are the stable corrosion products, Eq. (21.12) may be generalized [8] to yield p 3 c log ¼ a þ b logðPS2 Þ logðPO2 Þ þ ðtÞ 2 T

ð21:13Þ

Extensive exposure testing for many alloys and corrosive gases has shown that this form provides a good correlation of the corrosion data with the exposure conditions. For the specific case of Inconel 671 exposed to gas with the composition given in Table 21.1 for 900 K and 30-atm total pressure, the penetration is predicted by Eq. (21.13) to be 2.1 mils after 1 year, 4.6 mils after 5 years, and 6.5 mils after 10 years of exposure time [9]. For the same alloy under different service conditions (temperature, pressure, or gas composition), leading to a different combination of stable corrosion products, another empirical rate equation of a different form would be required.

D. SOFTWARE The foregoing approach to alloy selection is computationally intensive and the following logical steps are used: (a) The equilibrium partial pressures of elements in the corrosive gas must be found. (b) The stable corrosion product phases on the candidate metal or alloy must be predicted based upon (a). (c) A suitable empirical corrosion rate expression is selected, as appropriate for the likely corrosion mechanism [i.e., corrosion phases in (b)]. FIGURE 21.4. The Ni, Cr-H, O, S, C diagram at 900 K (627 C similar to Figure 21.3. The placement of the coordinate consistent with Table 21.1 shows that solid Ni3S2 is a stable phase coexisting with solid Cr2O3.

To assist in this manually cumbersome methodology, a user-friendly Windows 95 (trade name of Microsoft Corporation) program called ASSET (Alloy Selection System

286


for Elevated Temperatures) has been developed [9]. The thermochemical predictions involved in steps (a) and (b) are performed in a user transparent way with an appropriate subset of programming and databases of Facility for the Analysis of Chemical Thermodynamics (FACT) [2]. Regression analysis is used to produce correlations of alloy corrosion for four different corrosion mechanisms as named for the types of corrosion products that form: 1. 2. 3. 4.

Sulfidation Sulfidation/oxidation Oxidation Carburization

A large selection of data has been compiled for these conditions representing nearly 4.7 million h of exposure for 71 commercially available alloys. Details of the program are provided elsewhere [9].

E. CONCLUSION An engineering approach to the selection of alloys for hightemperature service in many different types of fossil fuel combustion and energy conversion processes, can, with suitable software [2, 9, 10], involve a blend of fundamental science and exposure testing. Thermodynamics is used to compute the equilibrium gas-phase chemical speciation and predict stable corrosion products on commercial alloys. Extensive exposure testing conducted on commercially available alloys is then correlated to those computations to provide a reliable guide to metal loss as a function of time, temperature, gas composition, and pressure. The reader is referred to readily available literature [8, 11–22] for a more comprehensive treatment of the corrosion of metals and alloys in high-temperature gases.

ACKNOWLEDGMENT The authors thank M. H. Kaye for his assistance in preparing the graphics.

REFERENCES 1. R. C. John, W. C. Fort, and R. A. Tait, Mater. High Temperatures, 11(1–4) (1993). 2. C. W. Bale, A. D. Pelton, and W. T. Thompson, “Facility for the Analysis of Chemical Thermodynamics,” in User Guide, Ecole Polytechnique de Montreal, Montreal, Canada, 1997. 3. G. Eriksson and W. T. Thompson, CALPHAD, 13(4), 377 (1989).

4. C. W. Bale, A. D. Pelton, and W. T. Thompson, Can. Metall. Q., 25(3) (July 1986). 5. M. Ahfat and W. T. Thompson, “Computation of Stoichiometric Phase Coexistence in Multicomponent Systems,” in Proceedings of the Metallurgical Society of CIM (Computer Software in Chemical and Extractive Metallurgy), Quebec City, Aug. 1993, pp. 99–111. 6. H. H. Uhlig and R. W. Revie, “Corrosion and Corrosion Control,” 4th ed., Wiley, New York, 2008. 7. N. Pilling and R. Bedworth, J. Inst. Met., 29, 534 (1923). 8. R. C. John, “Alloy Corrosion in Reducing Plus Sulfidizing Gases at 600–950 C,” in High Temperature Corrosion in Energy Systems, M. F. Rothman (Ed.), Metallurgical Society of AME, Warrendale, PA, 1985. 9. R. C. John, A. L. Young, and W. T. Thompson, “A Computer Program for Engineering Assessment of Alloy Corrosion in Complex High Temperature Gases,” in CORROSION 97, National Association of Corrosion Engineers, Houston, TX, 1997. 10. A. D. Pelton, W. T. Thompson, C. W. Bale, and G. Eriksson, High Temp. Sci., 26, 231 (1990). 11. J. D. Embury (Ed.), High Temperature Oxidation and Sulphidation Processes, Symposium by Metallurgical Society of CIM, 29th Annual Conference of Metallurgists, Pergamon, Elmsford, NY, 1990. 12. A. V. Ley (Ed.), Proceedings Corrosion-Erosion-Wear of Materials in Emerging Fossil Energy Systems, NACE, Houston, TX, 1982. 13. V. L. Hill and H. L. Blade (Eds.), The Properties and Performance of Materials in the Coal Gasification Environment, American Society for Metals, Metals Park, OH, 1981. 14. W. T. Bakker, S. Dapkunas, and W. Hill (Eds.), Materials for Coal Gasification, Conference proceedings of ASM Materials Week 1987, 1988. 15. G. Y. Lai, High Temperature Corrosion of Engineering Alloys, ASM International, Materials Park, OH, 1990. 16. P. Kofstad, High Temperature Corrosion, Elsevier Applied Science, New York, 1988. 17. R. Streiff, J. Stringer, R. Krutenat, and M. Caillet (Eds.), High Temperature Corrosion—Advanced Materials and Coatings, Vol. 2, Elsevier Science, New York, 1989. 18. R. A. Rapp (Ed.), High Temperature Corrosion, NACE, Houston, TX, 1983. 19. K. Natesan, P. Ganesan, and G. Lai (Eds.), Heat-Resistant Materials II, Conference Proceedings, ASM International, Materials Park, OH, Sept. 1995. 20. W. R. Davis (Ed.), Heat-Resistant Materials, ASM International, Materials Park, OH, 1997. 21. R. S. Treseder, R. Baboian, and C. G. Munger (Eds.), NACE Corrosion Engineer’s Reference Book, 2nd ed., NACE International, Houston, TX, 1991. 22. H. J. Gabke, Carburization—A High Temperature Corrosion Phenomenon, Materials Technology Institute, St. Louis, MO, 1998.

22 A PROCEDURE TO COMPUTE EQUILIBRIUM GAS-PHASE SPECIATION FOR USE WITH PREDOMINANCE DIAGRAMS M. H. PIRO AND B. J. LEWIS Department of Chemistry and Chemical Engineering, Royal Military College of Canada, Kingston, Ontario, Canada

W. T. THOMPSON Centre for Research in Computational Thermochemistry, Royal Military College of Canada, Kingston, Ontario, Canada

A. Introduction B. Homogeneous equilibrium in an ideal gas solution B1. Thermodynamic data B2. Identification of dominant species B3. Computation of equilibrium concentrations of gaseous species C. Heterogeneous equilibrium involving a pure condensed phase and an ideal gas D. Conclusion Acknowledgments References

A. INTRODUCTION In comparison to corrosion at near ambient temperatures, the oxidation of metals and alloys at high temperatures puts greater emphasis on chemical thermodynamics. Fortunately, the difficulties posed by experimental measurements at high temperatures are offset by more rapid chemical kinetics making computations of equilibrium especially useful as a predictive tool. In some cases, the decreasing stability of metal oxides with increasing temperature also makes exploiting conditions of immunity a practical possibility. A procedure is described in Chapter 21 of this Handbook to locate domains of immunity by computational means on isothermal diagrams in which the partial pressures of gases

are used as axes on logarithmic scales. These are sometimes called predominance diagrams and have a long history of usage in high-temperature corrosion as well as other matters involving chemical reactions at high temperatures between condensed phases and process gases, such as in smelting and refining. As the gas mixtures become more complex, particularly with respect to the number of elements, predominance diagrams become increasingly difficult to interpret. It is generally not evident how to create in a reacting gas mixture (such as CH4 and O2) the equilibrium partial pressures (CO, CO2, etc.) used as axes on predominance diagrams. Additionally, gas-phase equilibria controlling the speciation in the gas mixture are sensitive to temperature and pressure. Therefore, combinations of alloys and gas mixtures that perform well at one temperature may not do so at another. In these situations, multiphase equilibria states and speciation may be computed through numerical optimization techniques [1–5]. This method is applied to homogeneous gas-phase equilibrium in Section B and heterogeneous equilibrium of a condensed solid phase coexisting with an ideal gas in Section C.

B. HOMOGENEOUS EQUILIBRIUM IN AN IDEAL GAS SOLUTION High-temperature oxidation and corrosion requires the computation of gas-phase speciation at the temperature and pressure of exposure in a system comprised of several elements. To give focus to the computational methods, consider a hydrocarbon combustion situation. The usual


287

288

A PROCEDURE TO COMPUTE EQUILIBRIUM GAS-PHASE SPECIATION FOR USE WITH PREDOMINANCE DIAGRAMS

excess of oxygen in relation to a hydrocarbon fuel at high temperature makes it possible to exploit knowledge that the main product species are CO2(G) and H2O(G), with most of the sulfur appearing as SO2(G). However, with the same collection of elements involved in, for example, coal gasification, the deficiency of oxygen in relation to carbon and the large number of possible chemical compounds makes the computation of molecular concentrations more difficult. A general method to compute gas-phase equilibrium in a system with several elements is therefore a necessity. Even when the gas composition is known from chemical analysis at room temperature, the composition of the gas at the temperature and pressure of exposure must be computed. Gas-phase speciation is generally discussed with reference to equilibrium constants relating, for example, concentrations of CO(G), CO2 (G); and O2 (G): COðGÞ þ 12 O2ðGÞ ¼ CO2ðGÞ where

Keq ¼ exp

ð22:1Þ

DG PCO2 ¼ RT PCO PO20:5

ð22:2Þ

where Keq is the equilibrium constant, DG (Jmol 1) is the Gibbs energy of reaction, R (Jmol 1 K 1) is the universal gas constant, T (K) is the absolute temperature, and Pi (atm) is the partial pressure of species i in the gas. In an ideal gas mixture, the fugacity is numerically equivalent to the partial pressure in dimensionless units when the reference pressure is 1 atm. The difficulty with pursuing this method is incorporating the overall atomic proportion of carbon and oxygen in the system to express constraints on the partial pressures. Nonetheless, in a system where experience provides a basis for knowing the dominant gas species, and where the number of those species does not significantly exceed the number of

elements, this traditional method of computing gas-phase equilibrium is quite serviceable. However, it does not lend itself to generalization, even when applied to common combustion gases with several possible elements (C, H, S, O, N), particularly for the case of fuel consumed in a deficiency of air. B1. Thermodynamic Data As a basis of discussion that can be related to familiar expectations, consider the binary carbon–oxygen system at 1000 K and 1 atm. The Gibbs energy, Gi (Jmol 1), of an individual species i may be initially defined as [6] 0 Gi ¼ @DHi þ

ðT

1

0

Cp;i dT A T @Si þ

298

ðT

1 P ði Cp;i A dT þ Vi dP T

298

1

ð22:3Þ where DHi (Jmol 1) and Si (Jmol 1K 1) represent the standard enthalpy of formation and absolute entropy, respectively, at 298 K and 1 atm, Cp,i (Jmol 1K 1) is the molar heat capacity at constant pressure, Pi (atm) is the partial pressure, and Vi (m3mol1) is the molar volume occupied by i. Division by the total number of atoms in the formula mass (Mi,T) provides units of joules per gram-atom (Jg at 1). The use of this latter unit in expressing an extensive property provides a basis for comparing Gibbs energies of compounds with different numbers of atoms per molar mass on an equivalent basis. Although it is a misnomer, the Gibbs energy in Eq. (22.3), Gi , is sometimes called the “absolute” Gibbs energy. This nomenclature draws attention to the use of the absolute entropy in the formulation of Gi . The more usual Gibbs energy of isothermal formation from the most stable form of the elements appears on the extreme right side of Table 22.1.

TABLE 22.1. Gibbs Energies for Selected Species in C–O System at 1000 K and 1 atm Mole Fraction of Element

Atoms/ Molecule

“Absolute” Gibbs Energya [7]

Gibbs Energy of Formationa

Species

ci;C

ci;O

Mi;T

(Jg mol 1)

(Jg at 1)

(Jg mol 1)

(Jg at 1)

C(G) C2(G) C3(G) C4(G) C5(G) O(G) O2(G) O3(G) CO(G) C2O(G) CO2(G) C3O2(G) C(S)

1 1 1 1 1 0 0 0 0.5 0.667 0.333 0.6 1

0 0 0 0 0 1 1 1 0.5 0.333 0.667 0.4 0

1 2 3 4 5 1 2 3 2 3 3 5 1

547,992 618,794 563,277 711,512 698,392 77,353 220,764 119,445 323,264 29,161 629,423 409,605 12,660

547,992 309,397 187,759 177,878 139,678 77,353 110,382 39,815 161,632 9,720 209,808 81,921 12,660

559,952 644,114 601,255 762,151 761,690 187,735 0 211,702 200,222 164,863 396,000 150,862 0

559,952 322,057 200,418 190,538 152,338 187,735 0 70,567 100,111 54,954 132,000 30,172 0

a

In reference to iterations 0 and 1 in Table 22.2

HOMOGENEOUS EQUILIBRIUM IN AN IDEAL GAS SOLUTION

Taking CO(G) as an example, the species of pure carbon with the lowest “absolute” Gibbs energy (C(S)) and the lowest for pure oxygen (O2(G)) may be deducted from the “absolute” Gibbs energy of CO(G) to yield a Gibbs energy of formation of 100,111 (Jg at 1). Table 22.1 deliberately includes several species known to be of little consequence at 1000 K and 1 atm in order to illustrate the unbiased procedure by which the dominant chemical species are found as an initial step in computing the equilibrium composition of the gas phase. The atomic fractions of the elements in each species i are represented by ci;C and ci;O for carbon and oxygen respectively, as shown in Table 22.1 [7]. B2. Identification of Dominant Species The initial elemental proportions of carbon and oxygen must be specified to identify the dominant gas species. Here, this proportion is arbitrarily set to be 0.65 atom fraction O (equivalently, 0.35 mol C(S) and 0.325 mol O2(G)) for discussion purposes. This atomic proportion might be known from the relative flows of oxygen and pulverized carbon in a combustion situation. Temporarily setting aside the effect of concentration (or partial pressure) on the molar Gibbs energy of each species, the overall aim is to iteratively distribute the elements, subject to the constraint on their proportion, to make the Gibbs energy of the gas mixture the most negative. This approach could not be exploited in a practical way until the advent of digital computing using methods generally based on Lagrange multipliers [1–5] as Gibbs himself first recognized [8]. The present approach further exploits knowledge that Gibbs energy is, by nature, a relative quantity since it is not possible to chemically convert one element to another. Only changes in Gibbs energy, which involve processes of chemical conversion from one collection of C- and O-containing species to another, are measurable. Recognition of this thermodynamic principle permits systematic adjustment of the numerical values of the Gibbs energies for species in such a way as to preserve essential differences (as will be demonstrated later in Table 22.3). The process of adjusting Gibbs energies to preserve Gibbs energy differences for chemical change was coined by Eriksson and Thompson as “leveling” [9]. Leveling is performed by representing the set of Gibbs energies relative to the collection of species assumed to be most stable. The Gibbs energy of species i that is represented relative to a particular assemblage is referred to as the “relative Gibbs 1 ). energy,” DGm i (Jg-at The process is best understood using the sequence of diagrams in Figure 22.1 and the adjusted relative Gibbs energies in Table 22.2. Leveling involves: (i)

Selecting any two species which represent the overall atomic proportion of oxygen to carbon

289

(ii) Setting the relative Gibbs energy of those species to zero when they are assumed to be the most stable pair (iii) Readjusting the relative Gibbs energies of all other species relative to the two set to zero The two species initially selected are the most stable form of each element; that is to say, the form of carbon and oxygen with the lowest Gibbs energy in units of Jg at 1. The use of gram-atom in place of mole (or equivalently gram-mole) permits the direct comparison of equal amounts of monatomic O(G) with diatomic O2(G), as shown on the extreme right-hand side of Figure 22.1. With this initial selection of possible “dominant” species assures the ability to express the overall elemental proportion as represented by the vertical dotted line in Figure 22.1. All species are “leveled” with respect to C(S) and O2(G) in the first iteration, as represented by the horizontal dashed line at m ¼ 1 in Figure 22.1. The adjustment made to species i as a result of leveling (dGi) is represented by the following equation for the first iteration: dGi ¼ ci;C dGCðSÞ þ ci;O dGO2 ðGÞ

ð22:4Þ

By deducting this Gibbs energy evaluated at the composition of each C–O species, the relative Gibbs energies for all species are thereby set relative to C(S) and O2 (G) in the first iteration, and both elements in their most stable form necessarily become zero. This is depicted graphically at iteration m ¼ 1 in Figure 22.1 and tabulated in column 3 in Table 22.2. The next step involves selecting the species with the most negative relative Gibbs energy (i.e., CO2(G)), which will TABLE 22.2. Gibbs Energies Adjusted as Result of Leveling to Identify Dominant Species 1 Relative Gibbs Energy [DGm )] i (Jg at Iteration m

Species

0a

1a

2

3


547,992 309,397 187,759 177,878 139,678 77,353 110,382 39,815 161,632 9,720 209,808 81,921 12,660

560,652 322,057 200,418 190,538 152,338 187,735 0 70,567 100,111 54,954 132,000 30,172 0

560,652 322,057 200,418 190,538 152,338 385,735 198,000 268,567 1,111 120,954 0 49,028 0

565,097 326,502 204,863 194,983 156,783 383,513 195,777 266,345 0 123,177 0 50,806 4,445

a

In reference to Table 22.1.

290


Overall composition

Iteration m

250000

C3 (G) C4 (G) 125000 C 5 (G)

O(G)

C2O(G)

0

O3 (G)

C(S) -125000

C3O2 (G)

-250000 0.0

0.2

CO(G)

0.4

O2 (G)

CO2 (G)

0.6

0.8

-1 Relative Gibbs Energy, ΔG m i [J·gram-atom ]

250000

C3 (G) C4 (G) 125000 C 5 (G) 0

1.0

O(G) C2O(G)

O3 (G)

C3O2 (G)

O2 (G)

C(S)

-125000

CO(G)

-250000 0.0

0

0.2

0.4

1

CO2 (G) 0.6

0.8

1.0

250000

C3 (G) C4 (G) 125000 C 5 (G) 0

O2 (G)

C2O(G) C3O2 (G)

C(S)

CO(G)

2

CO2 (G)

-125000 -250000 0.0

0.2

0.4

0.6

0.8

250000

C3 (G) C4 (G) 125000 C 5 (G) C(S) 0

O2 (G)

C2O(G) C3O2 (G)

0.2

0.4

3

CO2 (G)

CO(G)

-125000 -250000 0.0

1.0

0.6

0.8

1.0

Atomic Fraction of Oxygen

FIGURE 22.1. Illustration of leveling technique applied to C–O system to determine pair of dominant gaseous species at 0.65 atomic fraction of O at 1000 K and 1 atm.

replace one of the species from the first iteration. The choice of the species to be replaced from iteration m ¼ 1 is made on the basis of being able to represent the fixed overall atomic fraction of O for the system as a whole. Thus, CO2 (G) must be paired with C(S) rather than O2 (G). A combination of CO2 (G) and C(S) are together more negative than the previous estimated assemblage comprised of C(S) and O2 (G). The equation representing the horizontal dashed line shown in Figure 22.1 is reestablished and leveling is advanced to iteration m ¼ 2 in Figure 22.1 and column 4 in Table 22.2. Species that have positive relative Gibbs energies do not require further consideration as they are correspondingly less chemically stable than the current assemblage. Thus, a large proportion of

possible species are withdrawn from further consideration with every additional iteration. This is a very significant feature when this method is applied to a gas mixture involving many elements and species. Eventually, all relative Gibbs energies are nonnegative, as in iteration m ¼ 3 in Figure 22.1 and the extreme right column in Table 22.2. The two species that have zero relative Gibbs energies at the end of leveling (CO2(G) and CO(G)) are the dominant gaseous species, and their proportions are as follows: 0:35 mol CðSÞ þ 0:325 mol O2ðGÞ ¼ 0:3 mol CO2ðGÞ þ 0:05 mol COðGÞ

ð22:5Þ

HOMOGENEOUS EQUILIBRIUM IN AN IDEAL GAS SOLUTION

Although the magnitude of Gibbs energies of all species changes throughout the leveling procedure, elemental differences in these energies are preserved at all iterations. The change in relative Gibbs energy for the reaction in Eq. (22.5) at iteration m is thus m m DGm Reaction ¼DGProducts DGReactants

DGm Reaction m ¼ 0:3mol 3 DGm þ0:05 mol 2 DG CO2ðGÞ COðGÞ m ¼ 0:35mol 1 DGm CðSÞ þ0:325 mol 2 DGO2ðGÞ ð22:6Þ The preservation of differences in Gibbs energy for the above reaction is demonstrated in Table 22.3 using relative Gibbs energies tabulated in Table 22.2 for all iterations. Differences in relative Gibbs energies of all other conceivable reactions in the system are also preserved. The thermochemical activity, am i (unitless), of species i is computed with respect to the relative Gibbs energy at iteration m through the following relation: am i ¼ exp

DGm i Mi;T RT

ð22:7Þ

The activity of species i is equivalent to its concentration, xm i;gas (and is correspondingly related to its partial pressure), in an ideal gas. Mi,T is the previously defined number of total atoms per formula mass. The estimated concentrations of all species are computed with Eq. (22.7) in Table 22.4 using relative Gibbs energies from iteration m ¼ 3. By the nature of this method, the mass constraints imposed on the system are in general not exactly satisfied after leveling. The overall mole fractions of carbon and oxygen in the gas are represented at iteration m by PN gas

m i¼1 Mi;C xi;gas PN gas m m i¼1 Mi;C xi;gas þ i¼1 Mi;O xi;gas

xm C;gas ¼ PN gas

TABLE 22.3. Iteration m 0 1 2 3

ð22:8Þ

291

and, PN gas

xm O;gas

m i¼1 Mi;O xi;gas PN gas m m i¼1 Mi;C xi;gas þ i¼1 Mi;O xi;gas

¼ PN gas

ð22:9Þ

where Mi,C and Mi,O are the number of carbon and oxygen atoms, respectively, in a molecule of i. The overall mole fraction x m gas representing the gas with estimated concentrations at iteration m will differ from the mole fraction that the gas is constrained by, x 0 gas . When the system is assumed to be homogeneous and comprised of a gas, the value for x0 gas representing the gas is equal to the overall mole fraction of the system, x* . The mass constraint x0 gas imposed on the gas is distinguished from the mass constraint x* imposed on the entire system for the condition where the system is heterogeneous and x0 gas 6¼ x . The treatment of heterogeneous systems involving an ideal gas coexisting with a pure condensed phase is addressed in Section C. m Discrepancies between x m C;gas and x O;gas with respect to the fixed mole fractions of carbon and oxygen represented by the gas (x 0 C;gas and x 0 O;gas ) are evaluated on a relative basis. The relative errors representing the overall mass balances are as follows: jm C

x 0 0 m m C;gas x C m x O;gas xO ¼ j ¼ x 0 C;gas O x0 O;gas

ð22:10Þ

By virtue of the dependency of concentrations of gaseous species on Gibbs energies, the sum of estimated concentrations in the gas, xm T;gas , will not equal unity. The relative error representing the deviation of xm jm x;gas T;gas from unity is thus defined m jm x;gas ¼ xT;gas 1

ð22:11Þ

The overall mole fraction xm gas represented by the gas and the corresponding relative errors are computed in Table 22.4. Table 22.4 demonstrates that the mass constraints imposed on the gas have not yet been satisfied at iteration m ¼ 3. The effect of concentration on the molar Gibbs energy, which was temporarily superceded during leveling, must now be

Elemental Differences in Gibbs Energies Preserved for All Iterations for Eq. (22.6) DGReactants (J)

DGProducts (J) (0.3)(3)( 209,808) þ (0.05)(2)( 161,632) (0.3)(3)( 132,000) þ (0.05)(2)( 100,111) (0.3)(3)(0) þ (0.05)(2)( 1,111) (0.3)(3)(0) þ (0.05)(2)(0)

(0.35)(1)( 12,660) þ (0.325)(2)( 110,382) (0.35)(1)(0) þ (0.325)(2)(0) (0.35)(1)(0) þ (0.325)(2)(198,000) (0.35)(1)(4,445) þ (0.325)(2)(195,777)

DGReaction (J) ¼ ¼ ¼ ¼

128,811 128,811 128,811 128,811

292


TABLE 22.4. Estimated Equilibrium Gas-Phase Composition Computed from Adjusted Relative Gibbs Energies from Iteration m ¼ 3 in Figure 22.1 and Table 22.2 Species

Mi;C

Mi;O

Mi;T

DGm¼3 (Jg at 1) i

am¼3 i

xm¼3 i;gas

Mi;C xm¼3 i;gas

Mi;O xm¼3 i;gas


1 2 3 4 5 0 0 0 1 2 1 3 1

0 0 0 0 0 1 2 3 1 1 2 2 0

1 2 3 4 5 1 2 3 2 3 3 5 1

564,397 326,502 204,863 194,983 156,783 383,513 195,777 266,345 0 123,177 0 50,806 4,445

3.30 10 30 7.75 10 35 7.87 10 33 1.82 10 41 1.12 10 41 9.26 10 21 3.52 10 21 1.82 10 42 1.00 4.98 10 20 1.00 5.38 10 14 5.86 10 1

3.30 10 30 7.75 10 35 7.87 10 33 1.82 10 41 1.12 10 41 9.26 10 21 3.52 10 21 1.82 10 42 1.00 4.98 10 20 1.00 5.38 10 14

3.30 10 30 1.55 10 34 2.36 10 32 7.26 10 41 5.62 10 41 0.00 0.00 0.00 1.00 9.96 10 20 1.00 1.61 10 13

0.00 0.00 0.00 0.00 0.00 9.26 10 21 7.04 10 21 5.47 10 42 1.00 4.98 10 20 2.00 1.08 10 13

2.00

Sum m¼3 xm¼3 C;gas , x O;gas m¼3 m¼3 jm¼3 x;gas , jC , jO

considered. This matter is the principal cause of the relative errors shown at the bottom right of Table 22.4 (i.e., jCm¼3 , m¼3 ). The identification of the dominant species by the jO leveling method provides good initial estimates for further computation. B3. Computation of Equilibrium Concentrations of Gaseous Species In many circumstances, depending on the temperature and pressure, the dominant species determined by the leveling algorithm are the only species that contribute significantly to the equilibrium population and all other species are of inconsequential proportions. However, further computation is often required when minor species contribute appreciably to the equilibrium composition in addition to the dominant species. As a continuation of the concept of leveling, which exploits the fact that Gibbs energy is a relative function, a method is employed that exploits the knowledge that the equilibrium concentration of a solute species dissolved in a solution is related to its Gibbs energy. The objective is to further adjust the Gibbs energies of the pure elements to satisfy mass constraints within an admissible relative error. However, any slight numerical perturbation in the Gibbs energy of the pure elements (i.e., dGC and dGO) may result in m large transformations in xm C;gas and x O;gas . To further complicate matters, changes made to the Gibbs energy of pure carbon will not only result in changes in the estimated number of moles of carbon but also in the number of moles of oxygen, since the dominant portion of molecules in the system are comprised of both carbon and oxygen atoms.

1.00

2.00

3.00

4.00 10 1

6.00 10 1

1.43 10 1

7.69 10 2

Detailed explanations of the methodologies employed by this technique are discussed by Piro [10]. In view of the nonlinear dependency of xm C;gas on changes to Gibbs energies of both carbon and oxygen, xm C;gas is differentiated with respect to changes in the elemental Gibbs energies of both carbon and oxygen. The resulting partial differentials represent the individual contributions to changes in xm C;gas contingent on changes in elemental Gibbs energies. The numerical problem is reduced to a system of nonlinear equations in Eq. (22.12), which provides a mathematical framework for solving for a set of Gibbs energy adjustments to satisfy mass constraints [10]: 2

qxm C;gas 6 qG C 6 6 m 6 qxO;gas 4 qGC

3 3 2 0 qx m x C;gas C;gas m x 6 xm C;gas 7 qGO 7 7 dGC 7 6 T;gas 7 7 6 ¼ m 0 7 7 6 qxO;gas 5 dGO 4 x O;gas xm 5 O;gas qGO xm T;gas ð22:12Þ

The partial differentials in Eq. (22.12) must be computed with respect to Eq. (22.7) in order to solve for the Gibbs energy adjustments for the succeeding iteration. Each term of the square matrix in Eq. (22.12) is computed with the following formulation: PN gas Mi;j Mi;k m xi;gas qx m i¼1 RT j;gas ¼ PN gas P N gas m m qGk i¼1 Mi; j xi;gas þ i¼1 Mi;k xi;gas

ð22:13Þ

HETEROGENEOUS EQUILIBRIUM INVOLVING A PURE CONDENSED PHASE AND AN IDEAL GAS

293

TABLE 22.5.

Estimated Concentrations of All Species at m ¼ 3 Used to Compute Coefficients of Eq. (22.12) Using Eq. (22.13)

Species

Mi;C

Mi;O

xm i;gas

qðMi;C xm i Þ qGC

qðMi;C xm i Þ qGO

qðMi;O xm i Þ qGC

qðMi;O xm i Þ qGO


1 2 3 4 5 0 0 0 1 2 1 3 1

0 0 0 0 0 1 2 3 1 1 2 2 0

3.04 10 30 7.79 10 35 7.90 10 33 1.83 10 41 1.13 10 41 9.28 10 21 3.53 10 21 1.83 10 42 1.00 4.99 10 20 1.00 5.39 10 14

3.66 10 34 3.75 10 38 8.55 10 36 3.51 10 44 3.40 10 44 0.00 0.00 0.00 1.20 10 4 2.40 10 23 1.20 10 4 5.83 10 17

0.00 0.00 0.00 0.00 0.00 0.00 0.00 0.00 1.20 10 4 1.20 10 23 2.41 10 4 3.89 10 17

0.00 0.00 0.00 0.00 0.00 0.00 0.00 0.00 1.20 10 4 1.20 10 23 2.41 10 4 3.89 10 17

0.00 0.00 0.00 0.00 0.00 1.12 10 24 1.70 10 24 1.99 10 45 1.20 10 4 6.00 10 24 4.81 10 4 2.59 10 17

2.00

2.41 10 4

3.61 10 4

3.61 10 4

6.01 10 4

Sum

Equation (22.13) demonstrates that the adjustments made to the Gibbs energies of the pure elements take into consideration the molecular composition of each species and its estimated concentration from the current iteration. For example, this method recognizes that one molecule of CO2 (G) contains one carbon atom and two oxygen atoms. The computation of elemental Gibbs energy adjustments incorporates knowledge of the estimated contribution of all gaseous species in the system simultaneously to the overall mass balances of all elements represented by the gas. Table 22.5 exemplifies the computation of individual terms in Eq. (22.12) and (22.13) using estimated concentrations of species from iteration m ¼ 3 in Table 22.4. The coefficients of Eq. (22.12) incorporate values from the bottom of Tables 22.4 and 22.5. Components of the vector on the extreme right side of Eq. (22.12) are computed with Eq. (22.8) and (22.9), using mass constraints applied to the gas and the estimated mole fractions of the pure elements at iteration m ¼ 3. Insertion of computed values from Tables 22.4 and 22.5 into Eq. (22.12) gives 3 2 2 3 2:41 10 4 3:61 10 4 0:35 0:40 7 6 2:00 þ 3:00 7 2:00 þ 3:00 7 dG 6 6 6 2:00 7 C 7 6 ¼ 6 7 6 3:61 10 4 6:01 10 4 7 dGO 4 0:65 5 5 4 0:60 2:00 2:00 þ 3:00 2:00 þ 3:00

The relative Gibbs energies of all species in the system are further altered using Eq. (22.4), analogous to the adjustments made in the leveling method discussed in Section B2. Table 22.6 summarizes the adjustments made to the relative Gibbs energies for all species and the subsequent changes to m m xm i;gas , x gas and j . The process is repeated until the maximum m value for j is less than the specified error tolerance. A value of 10 5 has been found acceptable as the error tolerance limit for the majority of thermochemical scenarios. The concentrations of all gaseous species at iteration m ¼ 8 on the extreme right column in Table 22.6 are representative of the equilibrium composition. The concentration of each gaseous species is equivalent to its partial pressure since the total applied hydrostatic pressure is 1 atm for this example. The current example has involved a simple two-element system comprised of a few chemical species for discussion purposes. This algorithm can be easily extended to systems of many elements and a correspondingly large number of species. Performance statistics are shown in Figure 22.2 for thermochemical computations of homogeneous systems comprised of 200 species and up to 40 elements using the methodology described. Computation times are (obviously) dependent on many different factors, such as hardware performance, and are thus normalized relative to the computation time of 10 elements.

ð22:14Þ The resulting system of linear equations can be solved with a standard linear equation solver to determine the adjustments made to the Gibbs energies of the pure elements in the next iteration: 0:225 4:81 10 5 7:22 10 5 dGC ¼ 0:275 7:22 10 5 1:20 10 4 dGO ð22:15Þ

C. HETEROGENEOUS EQUILIBRIUM INVOLVING A PURE CONDENSED PHASE AND AN IDEAL GAS The strategy discussed in Section B can be applied to heterogeneous systems involving various combinations of pure condensed phases in equilibrium with an ideal gas. The inclusion of pure condensed phases is illustrated through the

294 TABLE 22.6. Progression in Estimated Concentrations of Gaseous Species Following Table 22.4 Continued Until the Maximum Relative Error Is Below an Arbitrary Tolerance (10 5) Iteration m¼4 Species C(G) C2(G) C3(G) C4(G) C5(G) O(G) O2(G) O3(G) CO(G) C2O(G) CO2(G) C3O2(G) C(S) xm T;gas xm C;gas xm O;gas jm x;gas jm C jm O

DGm i

(Jg at

577,568 338,973 217,335 207,454 169,255 378,316 190,581 261,148 3,638 129,759 693 56,210 16,917

1

)

m¼5 DGm i

xm i;gas 6.79 10 31 3.88 10 36 8.78 10 35 4.53 10 44 6.25 10 45 1.73 10 20 1.23 10 20 1.20 10 41 4.17 10 1 4.64 10 21 7.79 10 1 2.09 10 15 1.20 3.77 10

(Jg at

589,086 350,491 228,853 218,972 180,773 372,469 184,733 255,300 6,473 135,488 634 60,782 28,435

1

)

m¼6 DGm i

xm i;gas 1.70 10 31 2.43 10 37 1.38 10 36 1.77 10 46 6.14 10 48 3.50 10 20 5.03 10 20 9.87 10 41 2.11 10 1 5.87 10 22 7.96 10 1 1.34 10 16 1.01

1

3.58 10

1

(Jg at

595,141 356,547 234,908 225,028 186,828 369,224 181,488 252,055 7,878 138,444 489 63,117 34,490

1

)

m¼7 DGm i

xm i;gas 8.20 10 32 5.66 10 38 1.55 10 37 9.64 10 48 1.61 10 49 5.18 10 20 1.10 10 19 3.18 10 40 1.50 10 1 2.02 10 22 8.38 10 1 3.28 10 17

(Jg at

596,177 357,583 235,944 226,064 187,864 368,626 180,890 251,458 8,097 138,935 436 63,499 35,526

1

)

m¼8 DGm i

xm i;gas 7.24 10 32 4.41 10 38 1.06 10 37 5.86 10 48 8.63 10 50 5.56 10 20 1.27 10 19 3.95 10 40 1.43 10 1 1.69 10 22 8.55 10 1 2.61 10 17

(Jg at

596,173 357,579 235,941 226,060 187,860 368,615 180,880 251,447 8,090 138,929 427 63,493 35,522

1

)

xm i;gas 7.24 10 32 4.41 10 38 1.07 10 37 5.87 10 48 8.65 10 50 5.57 10 20 1.27 10 19 3.96 10 40 1.43 10 1 1.70 10 22 8.57 10 1 2.62 10 17

9.89 10 1

9.97 10 1

1.00

1

3.50 10 1

3.50 10 1

3.51 10

6.23 10 1

6.42 10 1

6.49 10 1

6.50 10 1

6.50 10 1

1

3

2

3

4.99 10 6

1.96 10

6.36 10

1.14 10

2.86 10

7.76 10 2

2.39 10 2

3.23 10 3

5.39 10 5

1.27 10 7

2

2

3

5

6.85 10 8

4.18 10

1.28 10

1.74 10

2.90 10

HETEROGENEOUS EQUILIBRIUM INVOLVING A PURE CONDENSED PHASE AND AN IDEAL GAS

Normalized Computation Time

12

900 K and 30 atm are listed in Table 22.7 for all species of interest in this system. Leveling is performed using the same methodology described in Section B2 whereby all species are initially treated as pure separate phases. The adjustments made to the elemental Gibbs energies of the system can be interpreted graphically as a hyperplane in four-dimensional Euclidean space, analogous to the horizontal dashed line shown in Figure 22.1 for a binary system. The leveling algorithm determines that C(S), CO2(G), H2O(G), and H2S(G) are the dominant species, as shown in Table 22.8 for iteration m ¼ 4. Therefore, the system is now assumed to be heterogeneous, consisting of graphite (C(S)) coexisting with an ideal gas phase. The Gibbs energies of the system are further adjusted using Eq. (22.12) to satisfy the mass balances of the system, similar to the previous example for a homogeneous system. Adjustments made to the elemental Gibbs energies are constrained by the relative Gibbs energy of C(S) when it is assumed to precipitate. That is to say that the hyperplane representing adjustments made to the elemental Gibbs energies of the system intersects the point corresponding to the relative Gibbs energy of C(S). The hyperplane pivots about this point and the intercepts correspond to adjustments made to the Gibbs energies of the pure elements. The mass constraints imposed on the gas (x0 gas ) differ from that of the system (x* ) in regard to a nonzero number of estimated moles of C(S). The number of moles of C(S) is added

10 8 6 4 2 10

20

30

295

40

Number of Elements

FIGURE 22.2. Performance statistics of thermochemical computations are shown for homogeneous systems (200 ideal species) represented by up to 40 elements. Computation times are averaged by one million unique thermochemical computations by varying composition, temperature, and pressure.

following example applied to the (C, O, H, S) system at 900 K and 30 atm in reference to the example in Chapter 21. The temperature, pressure, and elemental composition reflects coke/coal/oil gasification and several other industrial processes involving corrosion phenomena. The “absolute” Gibbs energies and Gibbs energies of formation computed at

TABLE 22.7. Absolute Gibbs Energies and Gibbs Energies of Formation Computed from Thermodynamic Database for (C, O, H, S) System at 900 K and 30 atma Absolute Gibbs Energy [7] Species CO(G) CO2(G) H2(G) H2O(G) CH4(G) H2S(G) COS(G) H2CO(G) C2H4(G) C2H6(G) H2S2(G) CH3SH(G) S2(G) SO2(G) SO3(G) O(G) O2(G) C(S) a

(Jg mol

1

)

274,541 577,321 103,521 399,959 234,304 194,527 339,792 303,908 142,158 286,433 220,677 246,164 64,724 512,325 624,888 121,364 171,146 10,309

This example is made in reference to Table 21.1 in Chapter 21.

Gibbs Energy of Formation )

(Jg mol 1)

(Jg at 1)

137,270 192,440 51,760 133,320 46,861 64,842 113,264 75,977 23,693 35,804 55,169 41,027 32,362 170,775 156,222 121,364 85,573 10,309

178,659 395,866 0 210,866 16,953 58,644 211,548 104,505 85,501 44,747 52,433 3,549 0 308,817 335,808 206,937 0 0

89,329 131,955 0 70,289 3,391 19,548 70,516 26,126 14,250 5,593 13,108 592 0 102,939 83,952 206,937 0 0

(Jg at

1

296


TABLE 22.8. Equilibrium Compositions Computed for Initial Gas Composition 35 mol CO þ 25 mol H2 þ 20 mol CO2 þ 19 mol H2O þ 1 mol H2S at 900 K and 30 atma Iteration m ¼ 4 Species CO(G) CO2(G) H2 (G) H2O(G) CH4(G) H2S(G) COS(G) H2CO(G) C2H4(G) C2H6(G) H2S2(G) CH3SH(G) S2(G) SO2(G) SO3(G) O(G) O2(G) C(S)

Iteration m ¼ 10

(Jg at 1) DGm¼4 i

am¼4 i

xm¼4 i;gas

DGm¼10 (Jg at 1) i

am¼10 i

xm¼10 i;gas

Pm¼10 i;gas (atm)

9,637 0 6,466 0 1,782 0 10,699 26,590 18,561 10,443 12,981 12,521 45,711 44,253 75,926 404,870 197,933 0

1.22 10 1 1.00 2.43 10 1 1.00 3.77 10 1 1.00 2.99 10 2 8.90 10 6 5.15 10 6 1.08 10 4 3.42 10 3 2.71 10 4 4.56 10 5 4.96 10 7 3.78 10 15 5.94 10 20 1.59 10 19 1.00

1.22 10 1 1.00 2.43 10 1 1.00 3.77 10 1 1.00 2.99 10 2 8.90 10 6 5.15 10 6 1.08 10 4 3.42 10 3 2.71 10 4 4.56 10 5 4.96 10 7 3.78 10 15 5.94 10 20 1.59 10 19

11,209 2,097 8,980 2,725 3,794 10,822 20,893 28,632 20,236 12,328 27,956 18,770 73,147 55,494 85,140 407,991 201,066 0

5.00 10 2 4.31 10 1 9.07 10 2 3.35 10 1 7.93 10 2 1.31 10 2 2.30 10 4 2.25 10 7 8.97 10 8 1.89 10 6 3.23 10 7 2.91 10 7 3.23 10 9 2.17 10 10 1.71 10 20 2.09 10 24 4.57 10 24 1.00

5.00 10 2 4.31 10 1 9.07 10 2 3.35 10 1 7.93 10 2 1.31 10 2 2.30 10 4 2.25 10 7 8.97 10 8 1.89 10 6 3.23 10 7 2.91 10 7 3.23 10 9 2.17 10 10 1.71 10 20 2.09 10 24 4.57 10 24

1.50 1.29 101 2.72 1.01 101 2.38 3.92 10 1 6.91 10 3 6.76 10 6 2.69 10 6 5.66 10 5 9.70 10 6 8.72 10 6 9.69 10 8 6.52 10 9 5.14 10 19 6.27 10 23 1.37 10 22

nm CðSÞ

3.00 101

1.28 101

xm T;gas xm H;gas xm C;gas xm O;gas xm S;gas jm x;gas jm H jm C jm O jm S

3.78

1.00

5.12 10

1

3.96 10 1

1.30 10

1

1.86 10 1 4.14 10 1

2.69 8.85 10

2

4.40 10 3 4.97 10 6

2.78

7.26 10 6 5.79 10 6 1.49 10 6 1.78 10 6

a The leveling method (m ¼ 4) provides initial estimates for the computation of the final equilibrium state (m ¼ 10). The computation of this system is made in reference to Table 21.1 in Chapter 21.

or withdrawn from the system every iteration in a systematic manner. An initial estimate of the number of moles of the pure condensed phase can be computed from the mass balance of the dominant species produced by the leveling algorithm at m ¼ 4. The final equilibrium composition is shown on the extreme right column in Table 22.8, which is referenced to Table 21.1 in Chapter 21. A predominance diagram can be constructed using the methodologies discussed in this chapter by systematically adjusting the partial pressures of two components while maintaining constant temperature, pressure, and activities of the remaining elements. These diagrams are convenient methods to graphically identify the predominant corrosion products over a range of partial pressures of two gaseous species. Figure 22.3 reproduces a

predominance diagram from Chapter 21 for the system shown in Tables 22.7 and 22.8. The coordinates corresponding to the partial pressures of CO(G) and CO2(G) in the extreme right column of Table 22.8 determines that Ni3S2(S) is a stable phase coexisting with Cr2O3(S).

D. CONCLUSION The application of computational thermodynamics in determining the equilibrium gas composition and various condensed phases may be utilized for predictive characterization of system speciation and phase behavior. The methodologies described herein are particularly valuable as they do not rely on any prior knowledge of the coexistence of

REFERENCES

297

PH2O = 10.1 atm; PH2S = 0.392 atm Cr/(Cr + Ni) D/d—that is, when the chemical reaction takes place so rapidly at the solid–solvent interface that the solution is quickly saturated and remains so during the dissolution process. In this case, the dissolution rate, j is controlled by mass transport. Equation (28.1) reduces to Eq. (28.3): j¼

D ðCsat C1 Þ d

ð28:3Þ

The transport process is enhanced by convection due to density differences between bulk and saturated solution (natural convection) and/or by the

391

hydrodynamics of the system under forced convection. Expressions for the boundary layer d have been derived from first principles for both natural (or free) and forced convection for a variety of simple geometries, and Eq. (28.3) has been validated many times. This process is often called direct dissolution. 2. When K < D/d— that is, when the dissolution rate is phase boundary controlled as opposed to mass transport controlled. Equation (28.1) reduces to Eq. (28.4): j ¼ kðCsat C1 Þ

ð28:4Þ

The phase boundary reaction rate is then fixed by the movement of ions across the interface and hence is governed by molecular diffusion. The effective diffusion length over which mass is transported is proportional to (Dt)1/2, and therefore, the change in thickness of the specimen, proportional to the mass dissolved, varies with t1/2. This process is often called indirect dissolution. 3. When K ’ D/d, both phase boundary and mass transport are controlling. In this case of mixed control, the potential difference (Csat C1) can be seen as divided into two parts: Csat C is the part that drives the phase boundary condition, and C C0, the part that drives the transport process, so as to keep the dissolution rates for each process equal. The dissolution rate is extremely temperature sensitive, largely determined by the exponential temperature dependence of diffusion, and can be expressed by the equation

1 1 jT ¼ A exp B T T1 where A is the dissolution rate at temperature T1 and B is a model constant. In most practical dissolution problems, the reaction of a solid in a solvent leads to a multicomponent system; no longer is the chemical composition defined by a single concentration nor is there a single saturation composition at a given temperature; instead of pure dissolution the reaction is of the type S1 þ L1 ! L2 þ S2, the formation of S2 causing the interface composition to change. For porous solids such as refractories, with open porosity and with matrix materials being fine and highly reactive, both dissolution and penetration occur; hence, most slagging situations involve chemical attack of the matrix or lowmelting constituent phases, which disrupts the structure and allows the coarse-grained aggregates to be carried away by the slag movement. When penetration is more important than dissolution, another mechanism of degradation needs to be considered: structural spalling.

392

CORROSION OF REFRACTORIES AND CERAMICS

B1.3.3. Structural or Chemical Spalling. While spalling is a general term for the cracking or fracture caused by stresses produced inside a refractory, chemical or structural spalling is a direct consequence of corrosion penetration. It should not be confused with pure thermal or pure mechanical spalling. Structural spalling is a net result of a change in the texture of the refractory, leading to cracking at a plane of mismatch, at the interface of an altered structure and the unaffected material. When slag penetration does not cause direct loss of material, the slag does partially or completely encase a volume of refractory, reducing its apparent porosity (by sometimes more than one-half), causing differential expansion with the associated development of stresses. As a result, there is a degradation in material strength and stiffness, the appearance of microcracks, and eventually totally disruptive cracks parallel to the hot face of the lining; this is structural spalling. Quantification of the mass of spalled materials from the hot face has been attempted by Chen and Buyakozturk [13], who combined slag dissolution and spalling effects into one expression in terms of the residual thickness X and the rate of thickness decrease. They proposed expressions as a function of the hot-face temperature TH and the maximum depth of slag penetration DP, referred to the location of the hot face at time ti during the (i 1)th and ith spalling. The same authors have developed a very interesting notion to approximate the value of Dp using a critical temperature criterion; they postulated that, for a given system, slag penetration occurs only when temperature reaches a critical value Tc, which could be related to the temperature above which a certain percentage of liquid phase may still exist in the refractory (always under a thermal gradient). B1.3.4. Corrosion by Gases and Dusts versus Liquids. For corrosion by gases and dusts, the same qualitative relationships are applicable as for gas and liquid penetration into a refractory, except that wetting is not a factor. In gas corrosion, the driving force is the pressure gradient. Gas penetration is less rapid, in terms of mass, than that of wetting liquids; reactions do not necessarily commence at the hot face; and the depth of penetration of gases can be much greater than that of liquids. Moving down a thermal gradient, gases may condense locally to form a liquid, which then may initiate a dissolution–corrosion process and migrate still further into the solid; alternatively, gases may react chemically to form new compounds (liquid or solid). Once condensation has occurred, the basic criteria for liquid corrosion apply. The condensed liquid fills the pores and debonds the refractory material, creating thermal expansion mismatches, weakening, softening, swelling, and slabbing. Sulfur dioxide, usually from combustion or smelting operations, provides good examples of the condensation–

corrosion patterns experienced in practice. Condensation leads to the formation of sulfate in magnesia at 1100 C, in lime at 1400 C, and in alumina at 600 C. When alkali vapors (Na2O) are present with SO2, condensation of Na2SO4 overlaps with SO2 attack. Volatile chlorides also play a role. When dusts are carried out in hot gases, they may lead to deposition or abrasion. Deposition leads to scaling or caking with either beneficial or life-shortening consequences. Dusts entrained in gases that condense within refractories may facilitate the gradual adhesion of solid particles to the working surface. The resulting buildup can be either a hard scale or lightly sintered cake acting as a protective barrier. In less obvious cases, dust and condensing gases fuse together to create an invasive liquid. In other instances, scaling and caking of refractory walls may not be acceptable to a particular process, and periodic descaling operations can lead to degradation of refractories. Although no deposition may occur, abrasion can take place when refractory linings are bombarded by dusts. The extent of abrasive wear depends on the impinging particle size, shape, hardness, mass and velocity, angle of impact, fluid dynamics of the system, and viscosity of the eluant gas, among other parameters. Abrasion exacerbates corrosion. Exposure of a refractory either to a liquid or to its saturated vapor is thermodynamically the same. The differences lie in the transport mechanism. In approximate order of decreasing thermodynamic power, the oxidizing gases are NOX, Cl2, O2, HCl, CO2, H2O, and SO2, and the reducing gases are active metal vapors, NH3, H2, CxHy, CO, common metal vapors, and H2S. B2. Corrosion of Specific Classes of Refractories Detailed compilations of data on corrosion resistance of industrial refractories are not available because there are too many different products and too many variables to consider. Corrosion is also very seldom the only criterion to consider when selecting a material for a given application; other properties, such as thermal shock resistance and mechanical strength, must also be optimized, and there may be a need to compromise on corrosion resistance. To assess the corrosion resistance of different types of refractories, a rule of thumb is to consider the temperature limits, above which corrosion is considered to be excessive, established considering the expected durability. From one application to another, this durability does vary greatly as does the mode of degradation: corrosion–dissolution or penetration and structural spalling or both. In Table 28.1, two cases are distinguished, slagging resistance and hot gas corrosion resistance. The first case is represented by a reaction of the type S þ L, where L can be either a slag (from ironmaking, steelmaking, or copper smelting) or a flux (from other nonferrous metallurgy), either basic (CaO/SiO2 > 2) or acidic (Ca/SiO2 < 2, on a molar ratio), where S is a

CORROSION OF INDUSTRIAL REFRACTORIES

TABLE 28.1.

393

Guideline for Corrosion Resistance Based on the Criterion: Maximum Temperature Limit Not Reached for “Normal” Durability Temperature Limits ( C) of Slags and Fluxesa

Type of Refractories 1 Magnesia (M) Doloma (D) (M), (D) or (M þ D) þ Graphite Magnesia–chrome (MK) 2 High alumina (A) (A) þ Graphite 3 Clays þ (A) (65% Al2O3) Clays þ (A) þ C 4 Superduty fireclays (F) Medium–low duty (F) 5 High silica 6 Zircon–zirconia Silicon carbide a

Temperature Limits ( C) of Hot Gasesa

Basic

Acid

Oxidized

Reduced

1700 1700 1800 1700 1600 1700 1300 1550 1200 NR NR 1450–1600 2000 >1800 800 1800 1900 600 1450 600 1400 1200 1750 >2000 1200

1700 1700 1700 NR 1900 1700 NR 1700 NR NR 1600 1800 1600

NR ¼ not recommended.

given refractory with a given texture. The second case is represented by a reaction of the type S þ G, where G can be various hot gases, with or without dusts, either oxidizing (oxygen, air, or CO2) or reducing (CO, H2) (see B1.3.4). In Table 28.1, six types of refractories, from basic magnesia to acid silica and neutral zircon–zirconia, are considered. B2.1. Basic Refractories. Basic bricks are commonly used in the metallurgical industry: steelmaking, copper– nickel, and the cement industry. Magnesia is the most commonly used basic refractory, followed by dolomite. Magnesia bricks are made of various magnesia raw materials (sinters from natural magnesite, sinters from seawater magnesia, or grains of fused magnesia). Many other types of magnesia bricks are used, in combination with chrome (MK), alumina (NA), and doloma (MD). The most corrosion-resistant magnesia bricks are made of fused magnesia grains. Magnesia–chrome bricks were best bricks available before magnesia–carbon-bonded bricks, containing flake graphite additions, were developed for their slagging resistance. Dolomitic refractories are made of burnt natural dolomite, usually low in iron oxide (0.3%) and alumina (0.2%). As for magnesia, very efficient dolomite–carbon bricks have been developed. As shown in Table 28.1, dolomite performance may match that of magnesia for specific applications. B2.2. High-Alumina Refractories. High-alumina refractories are those that contain more than 65% Al2O3, composed of a-alumina (pure Al2O3 99.3 þ %) and mullite (Al2O3)3 (SiO2)2. For high corrosion resistance, electrofused alumina and mullite products should be considered. As for any type of refractory, there are many different classes of products— fired and unfired—with many different bond systems, a

commonly used one being a phosphate chemical bond system. Corrosion dissolution resistance is improved with increasing alumina content, but slag penetration resistance of such bricks then decreases, and hence, the structural spalling resistance also decreases. Improvements have been achieved with addition of chrome or silicon carbide, and, more importantly, of carbon. B2.3. High-Alumina Refractories with Clay. Highalumina bricks with clay refractories have alumina content between 45 and 65%; they may contain andalusite or mullite mixed with fire clays and binders. Corrosion resistance varies greatly with mineral composition and relatively low refractoriness (measured by the lowest temperature at which a liquid phase is present in the material). The maximum temperature limit, as indicated in Table 28.1, may not apply for specific applications. B2.4. Fireclay Refractories. Super duty, medium-duty, and low-duty fireclays are products containing alumina and silica, with alumina content varying between 40 and 45%, 30 and 40%, and 25 and 30%, respectively. They are used in a wide variety of kilns and furnaces because of their low price but are limited to low-temperature applications (1200–1250 C). B2.5. Silica Refractories. Silica bricks contain typically >93% SiO2 with minor amounts of lime (CaO < 3.5%), alumina (Al2O3 < 2%), and iron oxide (Fe2O3 < 1.5%). They have high corrosion resistance to acid slags but are sensitive to high temperatures and to thermal shock even at medium temperature levels. Bricks with reduced porosity and lower alumina content may be used to prevent rapid

394


degradation in the presence of alkalies, when hot gas corrosion prevails. B2.6. Zircon–Zirconia and Silicon Carbide Refractories. Among the many available zircon–zirconia refractories, zircon bricks, containing mainly zircon sand (ZrSiO4), have good wear resistance and poor wettability by molten metals at processing temperatures of up to 1450 C. Zircon is decomposed to zirconia and silica at 1540 C. The most common zircon–zirconia bricks are the very special electrofused A–Z–S bricks containing alumina–zirconia– silica made especially for glass-melting furnaces. Refractory materials with high zirconia content require the use of stabilized zirconia (CaO or MgO in solid solution) to avoid cracking on heating due to monoclinic to tetragonal transformation in the temperature range of 1000–1200 C, with a large volumetric change. Properly prepared zirconia shapes have very high corrosion resistance and are not wetted by metals, even by steels, with different oxygen content. Silicon carbide refractories are the nonoxide refractories most often used with carbon blocks, usually under reducing conditions. Many different types of silicon carbide refractories are available. They have the following advantages over oxide refractories: high thermal conductivity, superior spalling resistance, superior abrasion resistance, and hence high corrosion resistance against nonoxidizing slag (as in nonferrous metallurgical applications). Properties and costs of SiC refractories vary greatly, depending on the sintering aid used, but they do tolerate much more “impurities” than do structural SiC ceramics, which are discussed in the following sections.

C. CORROSION OF STRUCTURAL CERAMICS The main differences between industrial refractories and structural ceramics are their purity and texture. Structural ceramics are manufactured from purified synthetic aggregates rather than from as-mined natural minerals. Since the 1960s, the refractory industry has used the available

synthetic raw materials, so that the differences between the two groups, although still very significant, are much less than they used to be. Porosity is a very important factor that influences texture; porosity is < 2% for most structural ceramics, whereas it is 12% or more for most refractories (up to 20% for conventional refractories and 50–70% for insulating refractories). Porosity is a consequence of the manufacturing process route followed, in particular the grain size distribution selected and the sintering conditions. For structural ceramics, mainly very fine grain sizes (in the 0.1–10-mm range) are used to obtain very dense materials. For structural ceramics, the objective is to maximize mechanical properties, whereas for refractories it is to maximize thermal shock resistance and volumetric change, with, of course, good corrosion resistance for both classes of materials. The fundamental principles of corrosion of industrial refractories also apply to corrosion of structural ceramics, in both the S þ L and S þ G cases. Nevertheless, as porosity of structural ceramics is much lower, dissolution is more important than penetration. For the S þ L case, the corrosion resistance of structural ceramics to molten salts determines their usefulness as components in the chemical industry, such as pumps and heat exchangers. The S þ G type of reaction is important for structural ceramics used as kiln furniture in industrial furnaces and as components in gas turbine and diesel engines. C1. Corrosion by Molten Salts Comparative information is presented in Table 28.2 on the corrosion resistance to fused salts, alkalis, and low-temperature oxides for two different classes of structural ceramics: oxides and nonoxides. Values have been compiled from the combination of two main sources [14, 15] and refer to the chemical inertness of pure crystalline materials, with purity >99.5%. For this reason, the data in Table 28.2 are of limited generality. For example, it is to be expected that ceramics with less purity will be attacked more rapidly. The code A, B, C should be used only for comparing the materials considered. It is to be

TABLE 28.2. Corrosion Resistance of Structural Ceramics to Fused Salts, Alkalis, and Low-Melting Oxides Fused Salts–Alkalis–Low-Temperature Oxidesa Ceramics at Purity (>99.5%) Alumina Zirconia (stabilized) SiC (react. bond.) Si3N4 (RBSN) Si3N4 (HPSN) BN (HP) a

NaCl

NaCl þ KCl

KNO3

Na2CO3

Na2SO4

KOH

Na2O

V2O5

A1000

A400

B900

A800 C800 C800

A800

B900

A900 C900 C900 C750 (air) C900 (air) C900 (air)

A1000 A1000 C1000 (air) C1000 (air) B1000 (air) C1000 (air)

B500 B500 C500 (air) C500 C500 C500

B500 C500 C600 C500 B500 C500

C800 C800 C800 C800 C800

A400 A400 A400

A—Resistant to attack up to temperature indicated ( C). B—Some reaction at temperature indicated. C—Appreciable attack at temperature indicated.

CORROSION OF STRUCTURAL CERAMICS

assumed that the information is based on short-term observations at ambient pressure. McCauley [2] and Lay [3] provide a more complete coverage of the subject, including attack by glasses, aqueous solutions, and molten metals, but no further generalizations can be made from the specific studies that they reviewed because of limitations in the scope of the studies. The data in Table 28.2 show that both dense oxides and nonoxides are susceptible to attack at low and intermediate temperatures. The normally protective SiO2 layer that forms on SiC and Si3N4 may explain why such ceramics have a poor corrosion resistance in basic salts. C2. Corrosion by Hot Gases The S þ G reactions discussed below refer to oxidation at high temperatures in air and oxygen, gaseous corrosion in the presence of condensed deposits, and reduction or oxidation by hot gases, CO(g), H2(g), H2S(g), and H2O(g). Hightemperature oxidation in air or oxygen is the most important type of corrosion for carbide- and nitride-based ceramics. Although the corrosion of structural ceramics in hot gases is similar in some aspects to the oxidation of metals and the same theories are used to describe the processes, there are some important differrences: 1. Ceramics have a much higher oxidation resistance than do metals. 2. Because oxidation occurs at much higher temperatures, the corrosion products formed are liquids and gases, as well as solids, which can lead to either weight gains or losses, causing practical experimental difficulties of interpretation. 3. The textures of ceramics are not as homogeneous as those of metals, and so it is very difficult to determine the controlling mechanism of the oxidation process without considering both the processing characteristics and the intrinsic nature of the raw materials used. Gaseous corrosion in the presence of condensed phase deposits in combustion applications is the second most

TABLE 28.3.

important type of corrosion, with sodium sulfate and vanadate being the most corrosive deposits. Nonoxide ceramics are particularly sensitive to this type of environment. In Table 28.3, the corrosion resistance of dense structural ceramics to hot gases are presented using the same code as in Table 28.2. The oxidation resistance of nonoxide ceramics is clearly less than that of alumina and zirconia. C3. Corrosion of Specific Classes of Structural Ceramics C3.1. Alumina. Alumina is considered to be one of the most versatile materials that resist corrosive attack by acids, alkalis, molten metals, molten salts, as well as oxidizing and reducing gases, even at high temperatures. It is widely used in many high-temperature laboratory and pilot plant scale applications, such as crucibles, tubes, and special shapes. C3.2. Zirconia and Thoria. Thoria is the most thermodynamically stable oxide and the one with the highest refractoriness, with a melting point of about 3300 C. Unfortunately, it is not a readily available material, being used essentially in the atomic energy industrial sector. Zirconia is also a very stable pure oxide, but it has to be used only when stabilized, fully or partially, with CaO, MgO, or Y2O3 additions. Such additions make the solid solutions hZrO2–Xi more vulnerable to acid attacks and to alkalis. C3.3. Silicon Carbide. Silicon carbide has very good resistance to acidic solutions (even HF) and molten metals but is prone to attack under severe alkaline (basic) conditions, in particular in air. The SiC ceramics produced by reaction sintering and containing free silicon are more susceptible to oxidation and chemical corrosion than is single-phase SiC obtained by hot pressing. The SiC is very susceptible to attack by liquid sodium and potassium sulfate. On oxidation, SiC forms a layer of SiO2, but the possibility of forming SiO(g) increases with temperature and when the partial pressure of oxygen is decreased. The transition between “passive” oxidation (SiO2 protective layer) and “active” oxidation [SiO(g) without protection] occurs at low pressures

Corrosion Resistance of Structural Ceramics to Hot Gasesa

Ceramics at Purity (> 99.5%)

Air

Steam

CO

H2

Alumina Zirconia (stabilized) SiC (RB) Si3N4(RBSN) Si3N4 (HPSN) BN (HP)

A1700 A2400 A1200 B1200 B1250 C1200

A1700 C1800 B300 A220 A250 C250

A1700

A1700

A > 1000 A > 800 A > 900 A2000

A > 1000 — A > 800 A > 800

a

395

A, B, C, as in Table 28.2.

H2S

F(g)

A1000 A1000 A1000

A > 800 A > 1000

396


and high temperatures. SiC ceramics are more sensitive to hot corrosion than is Si3N4. C3.4. Silicon Nitride. Thermodynamic calculations of equilibria in the system Si3N4–O2 revealed that several reactions resulting in the formation of solids SiO2, Si2N2O, SiO, and gaseous N2, N2O, NO, SiO, and SiN are possible. Reactions leading to the formation of SiO2(s) are characterized by the most negative change in free enthalpy, but at high temperatures and low oxygen partial pressure, the possibility of reactions leading to the formation of SiC increases (significant at 1300 C and higher). The notion of passive and active Si3N4 oxidation has been well documented, but the controlling factor can be either diffusion of oxygen to the Si3N4–SiO2 interface, diffusion of nitrogen to the SiO2–air interface, or diffusion of impurity ions and sintering aids from the inner layers to the surface of the ceramic. A detailed survey of the subject is presented in [15]. For hot corrosion to complement the data listed in Table 28.2, only one significant result obtained on reactionbonding silicon nitride (RBSN) will be recalled. Although Si3N4 exhibits high corrosion resistance in the stream of combustion products of pure fuel (10 5% Na and V, 0.5% S) below 1400 C, it corrodes rapidly at 900 C when fuel contains 0.005% Na, 0.005% V, and 3% S. In liquid media, Si3N4 resists attack by acids up to 100–200 C. The Si3N4 exhibits higher resistance in alkaline solutions than in alkali melts and is rapidly attacked by (Na2SO4, Na2SO4–NaCl, Na2SO4 þ V2O5) melts but have good resistance to chloride melts NaCl þ KCl up to 600 C and up to 1100 C in NaF. C3.5. Boron Nitride. Boron nitride ceramics are usually prepared by hot pressing and normally contain a small proportion of boric oxide, a useful impurity that helps pressing but not its oxidation behavior at high temperatures. The BN ceramics are reported to be attacked by strong acids but are relatively inert toward alkalies. The nonwetting behavior of BN, similar in some ways to pure graphite, may account for this behavior. Boron nitride starts to oxidize at about 700 C. Oxidation mechanisms differ from Si3N4 and SiC because in this case the oxide layer first formed, B2O3, is liquid and tends to evaporate to BO(g), BO2(g), and B2O3(g) as temperature increases. Under reducing conditions, pure BN may be usable up to 1800 C. The most valuable characteristic of BN is its resistance to wetting by metals and alloys.

D. PREVENTING CORROSION Minimizing corrosion of refractories and ceramics under severe conditions at high temperature is a very important task that requires the following three sets of problems to be addressed simultaneously:

1. Material selection, considering both intrinsic and extrinsic characteristics, such as the design and configuration of the component 2. Installation methods and maintenance procedures, considering that each material is a part of a bigger identity 3. Process control to minimize variability and extreme values In this section, only the first set of problems will be considered, although, for practical solutions, the other two sets of problems should not be overlooked. To select the most appropriate material, the first rule is to make the acid–basic character of the refractory or ceramic constituents similar to that of the corrosive fluids (liquids and/or hot gases) and then to control the penetration–dissolution mechanisms to improve the corrosion resistance. The single, most important factor, in terms of material properties, is the porosity and in a broader sense the texture of the material. D1. Porosity and Texture Improvements in the corrosion resistance of refractories have been obtained through texture control resulting from the evolution in manufacturing processes toward much larger size distribution, better mixing, better pressing, and purer raw materials. Total open porosity has decreased to 12%, a level of porosity that is required to maintain good thermal shock characteristics and acceptable insulating properties. While maintaining the porosity at this value, further improvements are possible through reducing the pore size distribution, so that the larger pores are of a smaller size, and through reducing the permeability of the refractory by using fused grains instead of sintered grains to reduce the openness of the porosity. Other possible ways to minimize penetration in refractories will be discussed in subsequent sections. For structural ceramics where high strength is more important than thermal shock resistance, porosity can be controlled to very low values (IV

KZ>VI

KZ>VIII

0.5 1 Ratio of crystallographic radii

KZ>VIII

1.5

FIGURE 29.2. Comparison of crystallographic ratios of cation to anions with respective ratios for glasses.

cation–anion radii ratios and coordination numbers of various cations in crystalline structures and oxide glasses is given in Figure 29.2. Average crystallographic radii are obtained from bond length data of Shannon [31] (with ˚ [32] instead of 1.26 used by Shannon). roxygen ¼ 1.32 A Figure 29.2 shows a 1inear relation clearly indicating that bond length in glasses and crystals are similar, but coordination numbers appear to be slightly higher in crystalline structures. According to this diagram, network-forming cations occur up to a ratio of radii of 0.55. B1. Heterogeneity When glasses contain multiple phases, typically one phase is more soluble than the other and may control the corrosion resistance or leaching rate. Examples are B-rich borosilicate glasses with an easily leachable B-rich glass phase or soluble alkali molybdate phases formed in ill-defined nuclear waste glass compositions. In contrast, in silica-rich borosilicate glasses, phase separation can lead to an increase in glass durability if a durable Si-rich continuous phase is formed, which includes the B-rich phase in dispersed form [33]. Even in the absence of phase separations, microheterogeneities may occur. In 1921, a glass structure theory was developed [34] based on randomly oriented microcrystals similar to the recent “strained mixed cluster” model [35, 36]. Typical clusters are boroxol rings in B2O3–glass, P8, or P9 cages in some phosphate glasses or layered Ca octahedra in Ca silicate glass. Cluster boundaries are reactive centers.

C. GLASS CORROSION MECHANISMS The corrosion rate of glass is not an inherent material property but depends on glass composition, structure, and surface states as well as on environmental conditions, such as

the composition of the corrosive fluid (including pH), hydrodynamic conditions, temperature, stress, strain, vapor pressure, and so on. An overall mechanism of sequential and parallel partial reactions can be formulated for the glass–water reaction, probably applicable to most glass compositions. Different reaction rates and empirical rate laws as well as specific selective leaching properties and pH changes are interpreted as resulting from the different relative importance and consequences of the various partial reactions in the overall reaction scheme. The nature of the rate-limiting reaction in this overall scheme depends on both environmental constraints and glass composition and structure. The principal partial reactions in the glass–water interaction are glass network hydration, ion exchange, and dissolution of the glass network, resulting in changes of composition of corroding fluids and formation of altered surface layers. Additionally, transport processes, such as access of reactants (e.g., water molecules) and removal of products (e. g., dissolved glass network formers), may become rate limiting under certain conditions. Glass hydration and alkali–H þ ion exchange are reactions that occur in parallel with glass network dissolution [37–40], and these reactions are interrelated; for example, in silicate glasses, alkali exchange leads in most cases to pH increase in solution (but not in Mg-rich solution), which increases glass matrix dissolution rates. The aIkali–H þ ion exchange process is often (not always [41]) diffusion controlled, implying that the rate of selective alkali release from glasses is initially highest and decreases with the square root of time until it becomes equal to the rate of matrix dissolution. The process may be governed either by the electroneutrality coupled, diffusion coefficients of alkali ions, and H þ -bearing species in opposite directions, with the lower mobility of H þ -bearing species such as H3O þ ions [42, 43] normally being rate controlling.

402

CORROSION OF GLASS

Alternatively, the rate-limiting reaction may be the diffusion of water molecules into the glass network [44–46], that is, the hydration of the glass network. Based on findings that the pH response of glass electrodes is essentially determined by exchange equilibria of the outermost surface species with species in solution, it was suggested that ion exchange at the surface is driven by electrochemical potential differences between charged and uncharged surfaces species and the bulk glass [47]. However, the electrical coupling of H þ and Na þ migration in the surface layer leads to electrical fields much stronger than those between the surface and the glass [48]. Various models for the coupling of ion exchange and matrix dissolution can be found in the literature [49–51]. The relative importance of the ion exchange reaction depends on the diffusion coefficient D (m2/s) and the corrosion rate r (m/s). The D depends on glass and solution composition. For borosilicate glasses it can be shown that D is not related to the diffusion coefficient of Na in the dry glass [52]. The ion exchange process is strongly reduced with increasing alkali content of the solutions. The effect increases in the order Li þ < Na þ < K þ < Cs þ [53]. In the time between t ¼ 0 and t r2/D, the overall reaction is diffusion controlled [54]. Matrix dissolution becomes dominant for times t r2/D. The ratio D/r gives the average steady-state depth of either alkali diffusion or water diffusion in the glass and the product rt denotes the thickness of corroded glass. For silicate glasses, it has been shown [55] that the corrosion rate r is proportional to the thermodynamic stability of a hypothetical mechanical mixture of component oxides. This linear free-energy relationship has been confirmed for window glasses, medieval church glasses, natural

basalt glasses, and nuclear waste glasses [56], In general, the corrosion rate r is not constant but depends on solution composition. The pH and the concentration of matrix formers (e.g., dissolved silica) in solution are particularly important. Minimal corrosion rates are often observed at neutral pH [57, 58]. The pH dependence is related to the concentration of activated surface complexes [59]. Often, a decrease of corrosion rate r with time is observed. Initially, this decrease with time was interpreted as an effect of protective layers, but later it was shown, for silicate glasses, to be an effect caused by solution saturation, described in a simplified way as [60–62] r ¼ kþ

CSi; actual solution concd: 1 CSi; saturaton concd:

ð29:1Þ

with the saturation concentration of silicic acid and the forward rate constant k þ depending on pH, solution composition, temperature, and ionic strength. Saturation effects were also reported for some phosphate glasses [212]. With the approach of saturation, the affinity [63] of glass dissolution and thus the reaction rate decrease [64–66]. At saturation, the corrosion rate decreases by up to a factor of 1000, but corrosion continues, due to either formation of secondary phase or continued alkali release by the ion exchange reaction. The saturation effect may result in an increase of corrosion rate with increasing solution volume. The effect of surface area to volume ratio (S/V) on glass corrosion rates has been known for more than 20 years [67, 68] and is illustrated with experimental data for a nuclear waste glass in Figure 29.3 [71].

FIGURE 29.3. Dependence of the extent of glass corrosion [normalized elemental mass loss (NL)] on the ratio S/V: PNL–76/86 glass, 90 C, pure water [67]. Divided by the density of the glass (2.7 g/cm3) NL values can be translated into equivalent reaction layer thickness (in mm).

403

EFFECTS OF GLASS COMPOSITION ON CORROSION KINETICS

At the lowest S/V, mass loss per unit surface area, corrosion rates, and corrosion layer thickness are maximum values and independent of S/V. Two further reaction steps can be observed with higher S/V [69]: an intermediate step where the reaction rate decreases to a minimum (as shown in Fig. 29.3) and at very high S/V, where the reaction rate is augmented by secondary phase formation. Whether or not this last step occurs depends on glass composition, particularity on the Al content [70]. Two S/V effects should be distinguished: pH-related effects and saturation effects. Larger S/V values lead to faster saturation and often to pH increase (not always, e.g., not with phosphate glasses, or with silicate glasses in Mg-rich waters). Thus, in a nonbuffered system, the pH-dependent corrosion mechanism [68] varies with the S/V ratio. The saturationrelated S/V effect leads to a dependence of the extent of glass corrosion on the product t. S/V, as confirmed experimentally [71]; however, simultaneous occurrence of saturation effects and pH variations may complicate this simplistic view [72].

D. EFFECTS OF GLASS COMPOSITION ON CORROSION KINETICS The stability of the glass matrix against corrosion, ion exchange, and water diffusion as well as saturation concentrations and long-term corrosion mechanism, depend to a large extent on glass composition. The chemical nature and concentration of the glass components are both important. Depending on the interactions with other glass components, a given component may act as either a stabilizing or a destabilizing agent; it may enhance or decrease the relative importance of ion exchange when compared to matrix dissolution.

TABLE 29.1.

Sodium-trisilicate Corning 015 soda lime PPG soda lime R7T7 nuclear waste borosilicate Pyrex borosilicate Kimble R6 soda lime

Tektite a

See [73].

ln r ¼ k1

X

Xi DGohyd;i þ k2

ð29:2Þ

where k1 and k2 are empirical constants that are specific for a given corrosion test. For borosilicate and lead-rich glasses, deviations from linearity were interpreted as resulting from different structural roles of boron in the glass and were corrected [79]. If solution hydrolysis and complexation of

Corrosion Rates of Silicate Glasses at 90 C, pH7a Composition (wt %)

Obsidian

A comparison of the corrosion rates of important commercial silicate glasses at 90 C and buffered pH 7 is given in Table 29.1 [73]. The table shows that glass components do influence glass stability. However, since effects are not additive, it is difficult to predict the corrosion rates of multicomponent glasses from the corrosion rates of simple glasses. The effect of composition on stability is not an inherent glass property but depends on solution composition, the S/V ratio, and so on. The corrosion stability of a given glass must be experimentally determined if performance is to be predicted for specific service conditions. Nevertheless, some quantitative estimations of glass corrosion effects are possible and are described in the following paragraphs. There are both theoretical and empirical models describing the effect of glass composition on corrosion rates. Theoretical models are based on semiempirical linearfree-energy relationships (LFER). The most important model is based on studies of Paul [75] showing a relation between corrosion resistance and the hydration enthalpy of a mechanical mixture of component oxides and silicates. This model applies particularly to k þ in Eq. 29.1 [76]. A list of standard hydration enthalpies DGohyd;i is available [77, 78] and can be used to formulate the standard hydration enthalpy of the glass phase in order to calculate corrosion rates caused by solution hydrolysis using the equation

SiO2 (75), Na2O (25) SiO2(72), Na2O(22), CaO(6) SiO2(74), Na2O(13.3), CaO(8.3), Al2O3 (0.06), MgO(3.7)

SiO2(82), Na2O(4), B2O3(14), Al2O3 (2) SiO2(74), Na2O(12.9), CaO(8.3), Al2O3(1.8), MgO(4.5) SiO2(76), Na2O(3.8), CaO(0.5), Al2O3(13), K20(4.8), FeO(0.7) SiO2(73), Na2O(1.5), CaO(1.9), Al2O3(13), K2O(2.4), FeO(4.4)

Ea (kJ/mol)

Hydration Free Energy (kJ/mol)

References

344 5.7 0.8

84 68

18.5 19.6 13.8

73 73 73

0.4

70

0.2 0.02

54 79

8.3 13.1

73 73

0.0002

62

3.6

73

0.08

80

Linear Rate (mm/d)

58

74

404

CORROSION OF GLASS

dissolved glass constituents as well as surface complexation constants are taken into account, the corrosion rates of silicate glass fibers can be accurately described in the pH range 1–12 [80]. An alternative model is the structural bond strength model [193]. This model assumes that the corrosion resistance is controlled by the average bond strength V in the glass, which can be calculated from the enthalpy of formation of the component oxides, considering the role of a glass component as network former, modifier, or intermediate oxide. For more accurate prediction, empirical models are used that allow interpolation (not extrapolation) in an empirical compositional multidimensional space. These models are based on a multivariate fit of thousands of experimental data to a general equation [81].

E. SURFACE MODIFICATION After the first contact of a corrosive fluid with the glass, the surface becomes electrically charged, bonds resulting from the population of charge carriers (surface species: :SiO , :SiOM, :SiOH, :POH ) that belong both to the glass and to the solution. Surface species distribution is responsible for the response function of glass electrodes, and not ion exchange, as is assumed normally [82]. Often, a sequence of reaction layers can be identified, consisting of an inner diffusion layer, followed by a porous gel layer, and one or more laminar layers of precipitated reaction products. The dissolution front of the glass phase is located between the gel layer and the inner diffusion layer. In many cases, there is a smooth transition between the diffusion layer and the gel layer and both together are termed the “gel layer” without distinction [91]. The chemical and mineralogical composition of the assemblage of layers determines the retention capacity for radiotoxic or chemical toxical substances, the reflection properties, and so on. The principal scheme of these layers is given in Figure 29.4

In water-containing fluids, the initial surface modification is glass hydration without loss of the glassy state. If sufficient network bonds are hydrolyzed, a transformed layer [83] (gel) is formed with a clear phase boundary to the glassy phase. The transformed surface layers are porous [84], containing molecular water [85, 86] and allowing for high ionic mobility [87] as well as high water mobility. Molecular water may result from inward diffusion of water carrying H þ ions or may alternatively, result from the condensation reaction [88, 89], 2:SiOH ! H2O þ :SiOSi: bonds. The quantity of water entering the glass depends, among other parameters, on space and strain in the glass network [90]. Ion exchange results in a depletion of alkali ions in the hydrated glass. Potassium-rich silicate glasses show concentration profiles with smoothly decreasing K concentration from the bulk glass to the surface, whereas sodium-rich silicate glass show S-shaped profiles with almost constant Na concentrations between an inflection point and the surface [91]. This difference in shapes was explained by the different structures of Na and K silicate glasses [92], but this is unlikely, as in mixed Na–K lime glasses the same difference in respective concentration profiles was observed [93]. Conradt and Scholze [91] provided a convincing explanation for the different shapes: When ion exchange in the “gel layer” is more rapid than the formation of this layer, the S shape is found, whereas for comparable velocities of the two processes the potassium profile results. In natural tektite glasses [94] and obsidians [95] one has observed water diffusion without alkali release, clearly indicating that ion exchange is a secondary process to water diffusion. Alkali release is also suppressed if the alkali content of the solution is high [96]. The corroded glass may also dissolve entirely (as occurs with some potassium silicate glasses), but in many cases the glass surface becomes enriched at the surface with sparingly soluble glass constituents. This is particularly true if the glass contains large quantities of heavy metals. Depending on glass

FIGURE 29.4. Schematic representation of different surface regions of corroded silicate glasses.

EFFECT OF FLUID MEDIUM

composition and environmental constraints (temperature, pressure, solution composition, pH, flow rate, vapor pressure, contact time, etc.), amorphous (gel) for crystalline reaction products are formed. When compared to the original glass, the gel layer is depleted in network-modifying elements (Na, Ca, Li, Cs, etc.) but enriched in Si, Al, transition, and heavy metal contents (Fe, Ti, Nd, U, Zr, etc.) [97, 98]. Under acid conditions, even rare earth elements become depleted [99]. Metal ions from solution may become incorporated into the gel layer (e.g., Mg). The content of molecular water is significantly higher than the content of silanol groups [100] and water may occur phase separated in the structure of the layer [101]. The interface between the diffusion layer and the gel layer often has the character of a phase boundary. The gel layer is attached to the diffusion layer by siloxane bridges (:SiOSi:) bonds [136] but is often possible to mechanically remove the adhering gel from the glass surface. The gel layer is a rather stable alteration product whose dissolution rate may become smaller [102] or larger [103] than that of the glass. It is likely [102] that the solubility of the gel determines the saturation effect with respect to the glass [Eq. 29.1]. Released glass components accumulate in solution until solubility limits with respect to pure and mixed secondary phases are achieved. If the solubility is exceeded, precipates form at the outer surface or even within the gel layer. The composition of the precipitates is a strong function of chemical alteration conditions (e.g., pH, pCO2, redox potential). In principle, the composition can be predicted both for closed and open systems using models based on equilibrium thermodynamics [104] (geochemical modeling [105]). Successful modeling efforts are reported in the literature, particularly on basalt glass alteration [106–108] and nuclear waste glass alteration [109, 110]. The sequence of secondary phases during silicate glass corrosion is metal (hydrate) oxides ! clays ! zeolites ! SiO2. Typical oxides and oxyhydroxides are amorphous or crystalline iron oxyhydrates such as goethite or akaganeite, Tihydroxides, or aluminum hydroxides such as gibbsite. These phases are formed in the first hours of glass/water contact within a layer of few hundred nanometers [111–114]. Clays are among the most abundant crystalline alteration products of both natural glasses and man-made glasses [97, 115–117, 148]. Zeolites are formed mainly at alkaline pH (8–14). Typical zeolites formed are phillipsite [118] and chabazite. Analcime is formed particularly at temperatures between 150 and 250 C [12, 97, 98, 115, 119, 148]. Other crystalline alteration products include calcium–silicate hydrate (CSH) phases BaSO4, anhydrite, molybdates, phosphates (e.g., Apatite [120], LiPO4 [148]), silicates (e.g., uranyl silicates, Weeksite [121], Uranophane, and Haiweeite [122]), and layered double oxides like hydrotalcite [123]. The kinetics of zeolite formation can be described

405

by growth and dissolution reaction processes, closely resembling the Oswald step sequence [124]. With certain glass compositions, secondary-phase formation may also be detrimental to glass corrosion. This is particularly the case if formation of silica-rich zeolites prevents silica concentrations in solution to reach saturation values. Frequently, an increase of corrosion rates is observed once zeolites are formed [125].

F. PROPERTIES OF SURFACE LAYERS F1. Surface Layers as Protective Surface Films In order to explain the time-dependent decrease of corrosion rates, it has been assumed [126, 127] that the corrosion layers are protective, but, in many instances, corrosion rates decrease because of saturation effects in solution [131]. Mass balance calculations have shown that borosilicate glass dissolution continues even if tightly adhering surface layers are formed [128]. However, particularly in Mg-rich solutions, a protective effect was observed for borosilicate glass compositions [129], soda lime glasses [130], and natural tektites [131]. For soda lime glass, the addition of Al had a similar effect as Mg in decreasing corrosion rates. Protective effects of surface layers were also observed on bioactive glasses exposed to artificial plasma [132] and phosphatecontaining solutions [133]. This effect was attributed to formation of Ca–phosphate crystals (apatite) at the glass surface. The protective effect of the surface layer may be interpreted by local saturation phenomena, that is, by the formation of a diffusion barrier for dissolved H4SiO4(aq) [134, 135]. This implies a protective effect being particularly important if the forward reaction rate of the glass is high. This protective effect becomes unimportant if the bulk solution becomes saturated with respect to dissolved H4SiO4(aq) or if, because of tensile stress, fractures are formed in the surface layer [136].

G. EFFECT OF FLUID MEDIUM G1. Effect of Composition of Aqueous Solution G1.1. pH The most important parameter governing glass corrosion is probably the solution pH. Solution pH influences the reaction mechanism, the solubility of silica and metal cations, the degree of selectivity of the corrosion process. For example, with borosilicate glasses in acid media, selective extraction of glass modifiers and of B is observed, whereas under alkaline conditions congruent dissolution occurs. Dissolution rates under alkaline conditions increase significantly with pH [137]. If the boron concentration in the glass is sufficiently high to hinder the formation of a continuous

406

CORROSION OF GLASS

three-dimensional Si–O–Si network, smaller (Si–O–Si)n units may be formed and may be released by a percolation process, leading to congruent release even under acid conditions (pH2) [138]. For evaluating the effect of pH on glass dissolution, it is important to recognize the inverse effect of glass dissolution on pH. A self-accelerating failure mechanism of glass occurs if glass dissolution leads to increase or decrease of pH, and corrosion rates increase because of these pH changes. The interrelationship may be illustrated with some examples: 1. In a nonbuffered static test with limited solution volume, the alkali released from commercial Alcontaining soda–lime silicate glass leads to a rapid increase in pH, with increasing dissolution rates. Rates become significantly higher than those of Pyrex glass, the latter showing less increase in pH during corrosion. The situation is entirely different if the pH is buffered at 7. Then the soda–lime glass becomes more stable than Pyrex glass [73]. The slower increase of pH when dissolving Pyrex glass was explained with the smaller Na content of this glass. However, a factor probably even more important is the boron content of the glass. Dissolved boron can act as a very efficient buffer, fixing the pH at a value of 9. The high stability of borosilicate glasses may, at least in part, be attributed to this buffer effect. A key factor controlling pH changes during glass corrosion is the alkali–B ratio in the glass. If this ratio is smaller than or equal to unity, there will always be sufficient B in solution to buffer the pH. 2. Silicate glass dissolution in Mg-rich solutions leads to decreased pH and increased corrosion rate of Al- or Ferich glasses. This effect also occurs during the alteration of natural rhyolitic glasses in arid environments [139]. A particularly detrimental effect of the pH decrease is the mobilization of heavy metals. This effect was observed when actinides were released from borosilicate nuclear waste glasses [140]. 3. The phosphoric acid release from phosphate glasses, as well as the fluoride release from heavy metal fluoride glasses, in static tests leads to a rapid decrease in solution pH and to an increase in glass corrosion rates, as discussed in the following paragraphs. Often, these pH changes occur in standard tests, making data interpretation difficult. Studies on the effect of pH on glass corrosion must therefore be performed in buffer solutions at constant pH, in very high solution volumes, or at high water flow rates. The design of more stable glass compositions may focus more clearly on this self-accelerating mechanism; for example, constituents that provide a pH buffer effect can be added to the glass formulation.

G1.2. Effect of Special Ions in Solution Highly concentrated alkali salt solutions are known to decrease glass corrosion rates [141]. This effect has been explained by the suppression of the initial ion exchange processes. The corrosion resistance of alkali borosilicate glasses can be enhanced significantly by adsorption of certain ions from solution, such as Al3 þ , Ca2 þ , Mg2 þ , or Zn2 þ [142]. The effect may in some cases be explained by the formation of protective layers. Also, the solubility of the surface gel layer may be decreased by adsorption of these ions. If the saturation effect of glass dissolution is associated with the solubility of the gel, then the decrease of solubility should have a strong effect on glass corrosion rates. Adsorption of ions from solution was found to decrease in the order Ca > Ba > Ag > Mg [143]. G2. Corrosion by Vapor-Phase Hydration and the Atmosphere There are many similarities but also significant differences when comparing glass corrosion in liquid water with vaporphase hydration. In vapor test hydration, only a surface film of adsorbed water molecules exists on the glass surface and dissolved glass constituents cannot be transported away. The thickness of the water film depends on relative humidity (RH) as well as on the hygroscopy of the glass phase and hence, on glass composition. For some glasses, even 30% RH leads to attack of the glass surface. The quantity of water adsorbed increases with time and with the quantity of leached alkali. The humidity resistance increases in the order soda–lime bulb tumblers < soda–lime tubing < commercial plate glass, television panels < commercial window glass, soda–lime bottle glass, soda–lime tubing < commercial light bulb (outside) < alumino silicate, alkali borosilicate tubing < opal tableware, borosilicate tubing [144]. Due to the small thickness of the surface film, one may think of the vapor-phase hydration process as being similar to glass dissolution at high surface area–solution volume ratio (S/V). In such a thin water film on silicate glasses, the pH can rise rapidly. Once the pH is > 9, the corrosion of the glass network may occur, reaction products become deposited, and the glass surface loses its transparency and reflectance. Vertically stacked flat glass sheets may become unusable if stored under humid conditions. On the other hand, in the case of glass windows, the rain always washes these reaction products off, surface pH decreases, and glasses remain stable for longer periods of time. For corrosion of window glass in buildings under attack of atmospheric agents, four stages can be distinguished [145]: 1. Adsorption of water and CO2 associated with the formation of alkaline surface pH, leaching and formation of secondary minerals. This surface precipitate can easily be removed by cleaning with water.

BIODEGRADATION OF GLASSES

2. If humidity contact continues, surface precipitates become thicker (1 mm) and calcite forms and the surface becomes colored but can still be cleaned with water. 3. Further contact with humidity leads to strong leaching and enrichment of alkali and alkali earth elements as salt minerals at the surface and the silicate network becomes corroded by forming residual silica-rich structures. After cleaning with water, the surface remains modified, and effective cleaning requires acid treatment or even the use of hydrofluoric acid. 4. In the last stage the surface becomes rough and even a hydrofluoric acid cannot clean the surface. High S/V ratios are also attained by vapor-phase hydration of stacked window glasses in humid environments, for example, during transport in mountain or maritime regions [145]. On the other hand, at a lower relative humidity of 75% and 22 C it takes 70 days before significant dealkalization occurs [146]. Alteration phases (e.g., CaSiO3) [147] are rapidly formed due to high surface pH conditions. Also, hydrated carbonate minerals may form upon contact with carbon dioxide from the air. It is recommended to prevent stacked glass sheets to remain for a long time in humid environments, that interleaving materials be used, which may buffer potential pH rises, and that storage areas be well ventilated. In a corrosion test of borosilicate glass with steam at 200 C (steam generated 0.25 mL liquid water at the bottom of the reaction vessel), water condenses on the glass without dripping and an effective S/V of 4000 m 1 was estimated [148]. Higher water volumes led to condensation of water at the glass surface and dripping and thus recycled rinsing of the sample. After an induction period of a few days, corrosion rates insteam were found to be higher (linear rates of 2.5 and 21 mm/day for two similar glasses) than rates (after 30 days 0.01 and 1 mm/days) under hydrothermal conditions at S/ V ¼ 40 m 1 at the same temperature. Corrosion layers in steam were found to be richer in alkali contents and contained more silica–rich crystalline phases (zeolites and feldspars) than the corrosion layers formed under hydrothermal conditions. It was suggested that these phases drive the long-term glass corrosion rate at high S/V ratios caused by uptake of silica and prevention of saturation effects. Vapor-phase hydration was also studied with phosphate glasses. Here no diffusion control was observed but a linear rate law was observed associated with hydrolysis of the glass network [149].

G3. Hydrofluoric Acid Etching of glasses by HF is used in many technological and scientific fields, such as glass surface cleaning, glass

407

strengthening, and fission track edging. Wet chemical etching of silicate glasses in HF solutions has been reviewed [150]. Up to HF concentrations of 5 wt. %, the etching rate increases linearly with HF concentration, whereas higher HF concentrations lead to an over proportional increase in etching rates. The mechanism of dissolution is governed by adsorption of HF and HF2 species and catalytic action of H þ ions [150]. The initial reaction step is HF adsorption, followed by leaching of alkali and alkali earth. The rate-limiting step appears to be the breakage of siloxane bonds both at the outer surface and in the interior of the leached layer [153]. The reaction between HF solutions and silicate glasses was found in certain cases to be transport controlled with respect to the mass transfer of dissolved glass constituents in solution with a typical activation energy of 20–45 kJ/mol [150–152]. Initial corrosion rates in 10 wt % HF solutions were 100 mm/day at 25 C for fused silica [150]. Corrosion resistance in HF decreased with the addition of Na, Mg [152], Al, and Ca [150] as well as of P or As to the glass, whereas the etching rate decreased after B addition [150]. The relations between glass composition and etching rate are highly nonlinear [153]. The etching of multicomponent glasses by HF does not always lead to complete dissolution. In particular, if the glass contains alkali earth or rare earth elements, the precipitation of sparingly soluble fluoride phases can be observed [150]. G4. Strong Acids Corrosion of borosilicate glass has been tested in strong nitric acid up to a normality of 16. A maximum corrosion rate of 2500 mm/day (2-h test at 25 C) was observed in 6 N solution, and higher acid strength led to a decrease in reaction rates, attributed to a decrease in water activity [154].

H. BIODEGRADATION OF GLASSES In contact with bacteria, many glasses show increased corrosion. There is evidence for the existence of “silicate bacteria,” defined as chemolithoautrophic bacteria, which gain their energy by deterioration of Si–O bonds in silicates [155]. The effect also depends on the availability of nutrition [156]. Microorganisms, such as molds, are known to affect the stability of optical glasses [157] as well as medieval window glasses [158]. Experimental corrosion measurements of model glasses of medieval glass composition in the presence of fungi have shown 5–30 more glass corrosion (gel layer formation) than in the absence of fungi [159]. The effect was explained by a combination of an acid attack on the glass surface and a corriplexation reaction.

408

CORROSION OF GLASS

I. SPECIAL GLASS COMPOSITIONS AND APPLICATIONS I1. Alkali/Alkali Earth Silicate Glasses— Window Glasses Among the most important industrial glasses are soda-lime glasses, that is, a special type of alkali/alkali earth silicate glasses. Alkali ions in general decrease the stability of silicate glasses by decreasing the cross-link density of the three-dimensional glass network and by provision of ion exchange sites. In the pH range 1.4–10.9 the reaction of pure alkali silicate glasses with water is characterized by selective alkali release, following initially a square root of time rate law [160]. Divalent or trivalent cations stabilize alkali silicate glasses by formation of O–R–O bridges in the glass structure [161, 162] which hinder the water diffusion/ion exchange process. For a soda–lime glass with the composition SiO2 75 wt %, Na2O 15 wt %, CaO 10 wt % exposed to a buffered pH5.3 solution at 88 C, the diffusion coefficients for H and Na where 6 10 16 and 6 10 13 cm2/s, respectively [163], probably governed by diffusion through a transformed surface layer. Diffusion coefficients decreased by about a factor of 5 if an aliquot of 3 wt % of the Na2O content of the glass is replaced by K2O, emphasizing a mixed alkali effect [164]. An isotopic effect D/H, as well as 18 O update data, is consistent with rate control by breaking an HO bond (or DO bond) of indiffusing water molecules and subsequent hydrolysis of SiOSi bonds to 2 SiOH as necessary precursor reactions to Na/H ion exchange [165] Isotopic exchange reactions indicated condensation reactions according to the Scheme 2:SiOH ! :SiOSi: þ H2O as well as high mobility of water in the leached layer, 1000 times faster than the rate of leached layer growth. Hence, water diffusion was not rate limiting. Condensation of silanol groups in the leached surface layer was also evidenced by an increase in Q4 silica groups for a potash–lime glass, analyzed by (magic-angle-spinning nuclear magnetic resonance) 29 Si MAS–NMR [166]. While ion exchange and diffusion processes dominate the initial stages of the glass water reaction, glass network corrosion subsequently dominates. For a soda–lime composition for flat glass SiO2 72.0 wt %, Na2O 13.9 wt % MgO 4.0 wt %, CaO 8.3 wt % activation energies for forward glass network corrosion rates were found to be 66kJ/mol [167] at a neutral pH, while increase in activation energy is observed at lower and decrease at higher pH [191]. A primary factor controlling soda-lime glass network corrosion in nonbuffered water is a pH effect associated with the S/V ratio. Under static dissolution conditions the pH rises with the progress of glass corrosion. The higher the temperature or the S/V ratio, the faster this pH rise occurs. If pH > 9, a rapid increase of corrosion rates occurs. This does not happen if the pH is buffered at neutral values [168]. Then

slow corrosion rates of about 1 mm/day are observed at 90 C [191]. The rather poor resistance in alkaline solutions results in staining of window glasses at the rainexposed side of buildings if constructions allow for contact with alkaline runoff water from concrete or cement walls [169]. Introduction of nitrogen by adding Si3N4 to the glass melt increased durability of window glass considerably (factor of 3 in 2 N NaOH) [170]. For a soda–lime float glass, the composition of top and bottom surfaces are not identical, which is attributed to tin from the bath and Fe enrichment at the bottom. The tin layer can easily be dissolved in 1 M NaOH [171]. The tin-rich layer was found to protect the glass, resulting in about one order of magnitude lower glass surface leach rates [172] and less hydrogen penetration [173] when compared to the tin-free top surface. An addition of 2.5 wt % of tin to soda–lime glass increases glass stability [174]. I2. Alumosilicate Glasses Addition of Al to soda–lime glasses leads to another class of rather stable glasses, the alumosilicate glasses. The replacement of half of the CaO content of a soda–lime glass by Al2O3 (composition SiO2 75 wt %, Na2O 15 wt %, CaO 5 wt %, Al2O3 5wt %, leads to a decrease of diffusion coefficients for hydration by H-bearing species by a factor of 5 to a value of 10 16 cm2/s and of Na by a factor of 50 to a value of 10 14 cm2/s [163]. An increase of the Al–Si–ratio of silicate glasses increases the network corrosion resistance [175] by formation of Si–O–Al bridges and immobilization of a fraction of Na ions for reason of charge compensation in the vicinity of these bridges. The ratio of mobile to immobile Na ions in the glass (exchangeable and nonexchangeable Na) depends on the molar Na–Al ratio in the glass. With a molar ratio or Na/Al 1, an increase in the alkali content of the glass does not lead to an increased glass corrosion rate, because there are no nonbridging oxygen in the glass structure [176]. There are also some cases with no effects of Al on corrosion stability of the glasses [177]. The formation of Si–O–Al bridges stabilizes the glass only if the hydrolytic stability of this bond structure is high. This is the case at neutral to slightly alkaline pH but not at acid pH. Also in strongly alkaline solutions corrosion was observed to increase with increasing Al contents, attributed to the ease of hydrolysis of Al–O bonds under alkaline conditions and the gradual change from tetrahedral to octahedral coordination of Al ions by contact to alkaline solutions [178]. Hence, the stabilizing effect of Al depends also on solution chemistry. In addition, the effect of alkali earth elements on the corrosion resistance of alumosilicate glass depends on the chemistry of the aquatic medium. For example, the partial replacement of Si by alkali earth ions decreases the corrosion resistance of alkali aluminum silicate

SPECIAL GLASS COMPOSITIONS AND APPLICATIONS

glasses (18.5 Na2O–7.4–Al2O3–74.1 SiO2) at neutral pH and increases it at pH 10 [179]. This effect results from the reduced mobility of alkali earth ions at high pH values. In addition, aluminum plays an important role in controlling the solution concentration of silica. This is important under closed system alteration conditions, where saturation effects may decrease glass corrosion rates. Two cases may be distinguished: (1) control of Si solution concentration by the transformed surface of the dissolving glass, leading to faster slowdown of glass corrosion rates due, to saturation [180, 181], or (2) control of Si concentration by the formation of secondary phases resulting in a hindrance of the slowdown of reaction due to a hindrance of saturation at the glass–solution interface. For example, an Al-rich glass showed a much higher long-term corrosion rate than a glass with much less Al [182], This was due to the reduction of silica concentrations in solution due to the formation of analcime, a zeolite mineral, as secondary solid reaction product. I3. Borosilicate Glass Borosilicate glasses with 4–8% alkali oxide, 2–7% Al2O3, 0–5% alkali earth oxides, 70–80% SiO2, and 7–13% B2O3 are particularly water and temperature resistant. Typical samples are Duran or Pyrex glass. Zinc added to borosilicate glasses at concentrations up to 16 mol % leads to increased glass stability [183]. In aluminum silicate glasses without NBO (Na/Al 1), a replacement of up to 80% of Al2O3 by B2O3 leads to an increase in corrosion resistance. The corrosion resistance of borosilicate glass can be understood only on the basis of glass structure. The relation

409

between glass structure, glass composition, and glass stability is described by a model of Dell et al. [184]. Including new results for the ternary system [186] as well as for the binary alkali borate [185] and alkali silicate system, the results of the model are illustrated in Figure 29.5 Besides SiO4 tetrahedra (Qn), additional structural units of borosilicate glasses are anionic BO4 tetrahedra (N4 sites) and symmetric or asymmetric trigonal groups (N3s or N3a sites), with a distribution among these sites determined by the mol ratio R ¼ Na2O/B2O3 and K ¼ SiO2/B2O3. The occurrence of tetrahedral groups depends, on the presence of alkali. The maximum tendency for phase separation into Si-rich and B-rich glass phases [186] occurs at R ¼ 0.19. The increase of alkali content of the glass increases the solubility of borate groups in the silicate glass network [184] by forming borosilicate groups related to Danburite groups and Reedmergnerite (NaBSi4O10) units [184, 186]. At higher R values, not all Na ions are associated with tetrahedral borate groups but to Si–Q3 tetrahedras or with further alkali content to Si–Q2 sites. This leads to an increase of alkali mobility in the glass, resulting in a stronger tendency for ion exchange with contacting aqueous solutions. High-silica glasses can be fabricated based on phaseseparated glasses in the Na2O–B2O3–SiO2 system (Vycor process [187]). Here, glass corrosion is not detrimental to glass use but is a key design parameter because the boraterich phase must be dissolved rapidly. The microstructure of the glass determines the rate of dissolution. Rapid dissolution can, for example, be achieved by adding a small amount of high-valency cations such as V, P, or Mo, increasing the size of phase separations in the glass and the rate of borate phase dissolution [188, 189].

FIGURE 29.5. Ternary Na2O–B2O3–SiO2 Structure unit stability diagram and relation to glass durability and phase separation.

410

CORROSION OF GLASS

An industrially important, durable borosilicate glass is Pyrex glass, (SiO2 80 wt %, Na2O 4 wt %, Al2O3 2 wt % B2O3 14 wt %). As expected from Figure 29.5, Pyrex glass is phase separated into a nearly pure silica matrix with a dispersed borosilicate phase with spherical regions of 5–10 nm diameter [190]. Corrosion rates are probably dominated by the dissolution of the boron-rich phase. Corrosion rates were found to be invariant with pH between 0 and 4 and increased by about a factor of 10 until pH 9. At 95 C and pH 9 a linear forward corrosion rate of 0.6 mm/day was measured with pHindependent activation energies of about 60 kJ/mol [191]. At 250–260 C in a closed system, boron release continued with a linear rate of 4 mm/day, but Na and Si loss reach a temperature-dependent saturation state, with Na/Si ratios in solution equal to those in the glass [192]. Activation energies for hydration were only 17 kJ/mol, much lower than the activation energies for B release. This indicates that hydration and B release are independent processes, and hydration does not occur by filling the holes of phase-separated boronrich phases. With simple borosilicate glasses, as well as nuclear waste borosilicate glass compositions, the corrosion resistance is not a linear function of glass composition but shows strong nonlinear behavior. For example, the addition of 0–2 wt % SiO2 to a waste glass for West Valley (USA) led to only a small increase in corrosion resistance, but 3% addition decreased the corrosion rate by a factor of 10 (Fig. 29.6). A further increase in silica content had a negligible effect [193]. For glass compositions containing redox-sensitive elements, the corrosion resistance becomes a function of the redox conditions during glass production as well as of the redox-state of the aqueous solution. For example, FeIII acts much more as a stabilizing agent than does FeII (Fig. 29.7) because FeIII may take a similar role in the glass structure as a network-forming element, such as Al. Consequently, if such

FIGURE 29.6. Dependence of the extent of glass corrosion (release of boron) on the content of silica in the glass [212].

FIGURE 29.7. Corrosion results at 90 C in MCC-3-Test (Na release) of two American radioactive waste glasses as function of redox state (FeII/Fe(tot)) of the glass matrix [212].

glasses are melted under oxidizing conditions, they become more stable when compared to glass melting under reducing conditions [193]. I4. Nuclear Waste Borosilicate Glasses Borosilicate glass is the principal solid matrix for immobilizing high-level radioactive waste for deep geological disposal. Numerous borosilicate glass compositions have been suggested [194]. Compositions of all industrially produced European radioactive waste borosilicate glasses are located in Figure 29.5 in an optimum balance between (1) low tendency for ion exchange, (2) relatively high corrosion stability, and (3) absence of phase separation. Higher corrosion resistance could be achieved by higher Si content (more Q4 groups), but this would increase the tendency for production of phase-separated glasses. The resulting structure is a stable three-dimensional glass network with N3, N4, Q3, and Q4 with coordination cages of oxygen of various sizes and coordination, capable of hosting a large suite of different multivalent radioactive cations simultaneously. The glass corrodes slowly in groundwater and humid air and inevitably certain quantities of radionuclides are mobilized. Various reviews on nuclear waste glass performance have been published [13, 14]. Predicting long-term glass performance is difficult. Disposal sites are located in deep

SPECIAL GLASS COMPOSITIONS AND APPLICATIONS

rock with very slow groundwater flow rates. Hence, saturation effects become much more important than for other types of glass applications. Accumulation of dissolved silica and other glass constituents in solution leads in many (not all) cases to a significant decrease in the overall reaction rate [195–197]. Provided the accumulation of Si is not hampered (i.e., by formation of Si-rich crystals such as zeolites, by high water flow), a saturation state will be reached with reaction rates < 1/1000 of the initial rates [198]. This decrease in reactivity has been attributed to decreasing affinity for the rate-limiting dissolution reaction [198–200]. Protective layers were normally not formed. The fundamental disequilibrium between the bulk glass phase and the solution remains, that is, overall affinity remains appreciable and the glass/water reaction will continue as long as an aqueous phase is present. The mechanism of this “long-term reaction” is not yet fully understood. It has been suggested that the dissolution rate may either be driven by the precipitation of secondary phases [198, 201], formation of colloids [196], or resumption of rate control by water diffusion/ ion exchange [202]. Glass alteration may or may not be accompanied by transfer of glass constituents (including radionuclides) to a potentially mobile aqueous phase. Sparingly soluble glass constituents become incorporated in or sorbed on solid glass alteration products on the glass surfacer—the gel layer and a precipitated layer—but may also be sorbed on groundwater colloids. The formation of new secondary phases, such as silicates, molybdates, uranates, carbonates, and so on, establishes a new geochemical barrier for reimmobilization of radionuclides dissolved from the waste matrices. As an example, secondary clay minerals (saponite), powellite, and baryte phases formed on the surface of Cogema type HLWglass during corrosion in brine. The phases formed are normally not pure but are solid solutions of quite complicated composition. Solid-solution formation is a beneficial effect for retention of radionuclides. The clay mineral saponite is important for sorption of trivalent and hexavalent actinides, while powellite and barite are host phases, incorporating either trivalent actinides or divalent radionuclides within their crystal structure. Tetravalent elements appear to become concentrated into thermodynamically very stable zircon or cereanite-type phases. Recently, it has been shown that the retention behavior of trivalent, tetravalent, and hexavalent elements can be described by sorption isotherms in certain pH ranges, whereas solubility/coprecipitation is dominant in other ranges [203]. Radiation damage of the glass phase and transmutation have only a minor effect on glass durability, whereas the effects of radiolysis of water and aqueous species may increase in glass dissolution rate by a factor of 5 [14]. Initial reaction rates may become higher if neutral starting solutions are acidified by radiolysis. Alternatively, for experiments at very high S/V in the presence of nitrogen, it has been

411

shown [204] that dissolution rates of radioactive glasses may be as much as a factor of 40 lower than the corresponding rates of a simulated glass. In this case, the alkaline evolution of solution pH is partly balanced by acidification due to radiolysis. In the absence of nitrogen the initial rate is slightly affected by radiolysis, but there is only little effect on the rate of long-term dissolution at high S/V [205, 206]. I5. Alkali-Resistant Glasses Glass-fiber-reinforced cement (GFRC or GRC) composites are based on glass fibers with high alkaline durability. Mechanical properties of these cements depend on the chemical durability of the fibers. Suitable alkali-resistant fibers contain high contents of Zr, for example, in CemFIL (16 wt %) or ARG-fiber. The Zr-containing glass fibers (CemFIL-1) were leached at 55 C for a year in Portland cement pore fluids, resulting in only 2 mm of corrosion. Glass stability was caused by the formation of a shell of Zr- and Ca-rich reaction products [207]. In the absence of Ca, no protective effects of Zr-rich surface reaction products were observed. Maximum glass stability was observed with 100 mg/L Ca in solution [208]. Disadvantages of the Zr-rich fibers are the expensive highmelting and processing temperatures. Therefore, a number of alternative alkali-resistant glass compositions have been developed. Glass compositions in the mole percent range CaO < 9, Na2O 13–15, BaO < 10, ZrO2 7, MgO 3, SiO2 65–75 show similar resistance to ChemFIL and 100 C lower manufacturing temperatures [209]. Low corrosion rates, 0.01 mm/d, of Zn–Al–silicate glass fibers, modified with Fe, Mn, or Ti oxides, were observed at 60 C in Ca(OH)2 solutions at pH 12.6 [210]. I6. Phosphate Glass Phosphate glasses show a number of advantages when compared to silicate glasses: optical properties, high thermal expansion coefficient, low melting points, and formation of sparingly soluble host phases for radiotoxic elements. However, instability to crystallization and low chemical durability of many phosphate glass composition lead to less widespread use than silicate glasses. Depending on glass composition, the corrosion rates of phosphate glasses cover a range between 1000 and 0.01 mm/day [211]. A particularly stable class of glasses is to be found in the lead–iron phosphate system. A review is given by Sales and Boatner [212]. I7. Glass-Reinforced Plastics Low specific gravity ( poly(3-hydroxybutyrate-co45% 3-hydroxybutyrate [P(3HB-co-45% 3HV)] > [P(3HBco-71% 3HV)] [121]. Similarly, the sequence of enzymatic hydrolysis follows [P(HB-co-16% HV)] > [P(HB-co-32% HV)] > PHB [169]. The crystallinity of the polymer also affects the rate of degradation but is rarely taken into account [119]. Structural substitution groups, and their number per repeating unit, also affect the degradation kinetics. For example, (CAs) with a lower degree of substitution (DS) values are more quickly degraded than higher substituted ones under both aerobic and anaerobic conditions [38–45, 120]. During degradation of CAs, both molecular weight and degree of substitution showed a decreasing trend, suggesting that deacetylation and decomposition of the polymer backbone proceed simultaneously [44]. Earlier data suggested that CAs with DS values greater than 0.82 are recalcitrant to biodegradation and the limiting step is deacetylation followed by breaking of the polymer carbon–carbon bonds [100]. The general theory of polymer degradation assumes that exoenzymes from microorganisms break complex polymers to yield short units, including oligomers, dimers, and monomers, small enough to be allowed to pass through permeable outer bacterial membranes and subsequently to be assimilated as carbon and energy sources. The process is called depolymerization. The closer a polymer structure is to a natural analog, the easier it is degraded. Cellulose, chitin, pullusan, and PHB are all biologically synthesized and are completely and rapidly biodegradable [170]. The complete decomposition of a polymer to CO2 and H2O under aerobic conditions or organic acids, CO2, and CH4 under anaerobic conditions is rare. Degradation and mineralization of a polymer substrate can hardly achieve 100% due to the synthesis of microbial biomass and humus matter [171]. Under different conditions, predominant groups of microorganisms differ and degradation pathways vary. In the presence of oxygen, aerobic microorganisms are responsible for destruction of complex materials, with microbial biomass, CO2, and H2O as the final products (Fig. 30.1). In contrast, under anoxic conditions, anaerobic consortia of microorganisms are involved in polymer deterioration, and the primary products will be microbial biomass, CO2, CH4, and H2O [142–144] (Fig. 30.1). These conditions are widely

426

MICROBIOLOGICAL DEGRADATION OF POLYMERIC MATERIALS

Polymer

Depolymerases Oligomers Dimers Monomers Aerobic Microbial Biomass CO2

Anaerobic Microbial Biomass CH4/H2S CO2

FIGURE 30.1. Schematic diagram showing degradation of polymeric materials under aerobic and anaerobic conditions.

found in natural environments and can be simulated in the laboratory with appropriate inocula. E1.1. Poly(b-hydroxyalkanoates). Bacterial poly(b-hydroxyalkanoates) are a class of aliphatic polyesters consisting of homo or copolymers of [R]- b-hydroxyalkanoic acids. The polymer is a microbial intracellular inclusion in the cytoplasmic fluid in the form of granules with diameters between 0.3 and 1.0 mm. They can be as much as 30–80% of cellular mass. Unlike other biopolymers, such as polysaccharides, proteins, and DNAs, PHB is a thermoplastic with a melting temperature around 180 C, making it a good candidate for thermoprocessing. Furthermore, a copolyester consisting of 3-hydroxybutyrate and 3-hydroxyvalerate, poly(3HB-co3HV) can be synthesized in cells of Alcaligenes eutrophus [172]. Both long side-chain and novel side-chain polymers have been synthesized by the addition of appropriate substrates in the culture medium and with different bacterial strains [173–175]. The copolymers range from thermoplastics to elastomers. The A. eutrophus genes encoding PHB synthesis have been transferred into cotton for PHB synthesis in fiber cells for production by agriculture [176]. Homopolymers and copolymers can be degraded in biologically active environments, for example, soils [91, 177], sludge, compost [38, 39, 43], river water [86, 92, 93], and seawater [86, 88, 89]. Extracellular PHB depolymerases have been isolated from P. lemoignei [178] and A. faecalis [140, 179]. The enzymatic degradation occurred at the surfaces of the polyester film, and the rate of surface erosion was strongly dependent on both the molecular weight (degree of polymerization) and composition of the polyester crystallinity, and the dominant species of bacteria. The stereochemistry of these chemically synthesized polymers also affects their fate after disposal in the natural environment [180, 181]. E1.2. Cellulose Acetates. Cellular acetates are chemically modified natural polymers designed to improve their

mechanical properties for different uses. Generally, CAs with a degree of substitution from 1.7 to 2.5 can be degraded in thermophilic compost in 45 days [38–45] and transformed into constituent chemicals through biologically catalyzed reactions [182]. Increasing the DS value on a repeating unit makes the CA less degradable. CA degradation occurs more rapidly under aerobic conditions. The mechanisms of degradation are deacetylation, which releases the substitution groups, followed by cleavage of the carbon–carbon backbone. The decrease in molecular weight and deacetylation proceed simutaneously during degradation. E1.3. Polyethers. Polyethers include polyethylene glycols (PEGs), polypropylene glycols (PPGs), and polytetramethylene glycols (PTMGs). They are widely used in pharmaceuticals, cosmetics, lubricants, inks, and surfactants. They frequently contaminate natural water, including coastal waters and streams where wastewater is discharged. Degradability of this class of polymers has been studied under both aerobic [162, 183–186] and anaerobic [187–189] conditions. Polyether degradability is dependent on molecular weight, with molecular weights higher than 1000 considered resistant to biodegradation [162, 190]. However, degradation of PEGs with molecular weights up to 20,000 has been reported. The ability of a microflora to degrade larger PEG molecules is dependent primarily on the ability of a syntrophic association in mixed cultures of bacteria to metabolize the chemicals. For example, Flavobacterium and Pseudomonas can degrade PEG. After each oxidation cycle, PEG molecules are reduced by a glycol unit. The central theme of PEG degradation is cleavage of an aliphatic ether linkage. In a coculture of aerobic Flavobacterium and Pseudomonas species, PEG degradation proceeds through dehydroxylation to form an aldehyde and a further dehydrogenation to a carboxylic acid derivative [184, 185]. Neither of these bacteria can degrade PEG in pure culture. Cellular contact between them seems to be essential for effective cooperation. In the Flavobaterium and Psuedomonas system, three enzymes are involved in the complete degradation of PEG [162]. PEG dehydrogenase, PEG–aldehyde dehydrogenase, and PEG–carboxylate dehydrogenase (ether cleaving) are required. All three were found in Flavobaterium, while only PEG–carboxylate dehydrogenase was present in Pseudomonas. However, Pseudomonas, though not directly involved in the degradation, utilizes a toxic metabolite that inhibits the activity of the Flavobacterium. This appears to be the essential link for their syntrophic association in the degradation of PEG. E2. Effects on Physical Properties E2.1. Polyimides. The wide acceptance of polyimides in the electronics industry [191–195] has drawn great attention to the issue of stability of these materials. The National

MECHANISMS OF DEGRADATION

Research Council in 1987 [3] emphasized the need to develop deterioration preventive measures for polymers used in the electronic industries because data acquisition, information processing, and communication are critically dependent on materials performance and integrity. The interlayering of polyimides and electronics in integrated circuits prompted several studies on the interactions between these two materials [196, 197]. Polyimides are also widely used in load-bearing applications, for example, struts, chasses, and brackets in automotive and aircraft structures, due to their flexibility and compressive strength. They are used in appliance construction, cookware, and food packaging because of their chemical resistance to oils, greases, and fats; microwave transparency; and thermal resistance. Their electrical properties are ideally suited for applications in the electrical and electronics markets, especially as high-temperature insulation materials and passivation layers in the fabrication of integrated circuits and flexible circuitry. In addition, the flammability resistance of this class of polymers may provide a halogen-free flameretardant material for aircraft interiors, furnishings, and wire insulation. Other possible uses may include fibers for protective clothing, advanced composite structures, adhesives, insulation tapes, foam, and optics operating at high temperatures [195]. Electronic packaging polyimides are particularly useful because of their outstanding performance and engineering properties. However, they are susceptible to degradation by the colonization of fungi (Fig. 30.2, Table 30.1)[25–31, 33– 35, 47]. Polyimide degradation occurs through biofilm formation and subsequent physical changes in the polymer. Using electrochemical impedance spectroscopy (EIS) [198, 199], fungal growth on polyimides yields distinctive EIS spectra, indicative of failing resistivity [20, 21, 26]. Two steps are involved during degradation: An initial decline in coating resistance is related to the partial ingress of water and ionic

FIGURE 30.2. Scanning electron micrograph of fungi on deteriorated polyimides (scale bar, 5 mm).

427

TABLE 30.1. Polymeric Materials Tested for Their Susceptibility to Degradation and Deterioration by Environmental Microorganisms Name

Description

References

Adhesive

RTV142 silicone rubber with methyl alcohol Benzophenonetetracarboxylic imide polymert foam Polytetrafluoroethylene, fluorinated ethylene propylene coated polyimides and perfluocarboxyl Fluorinated polyimide/glass fibers, bismaleimide/carbon fibers, epoxy/ carbon fibers unidirectional, epoxy/carbon fibers [0,45, 90, -45]2S, poly(ether-etherketone) Epoxy/graphite fiber unidirectional Epoxy/carbon fibers, epoxy/glass fibers, bismaleimide/aluminum Aliphatic polyurethane coating Pyromellitic dianhydride and 4,40 -diaminodiphenyl ether

[28]

Insulation foam Cable insulation

Composites

Protective coating Kapton polymimdes

[28] [28]

[28, 34, 35, 36, 200]

[28] [28, 36, 200]

[28] [26, 30–32]

species into the polymer matrices. This is followed by further deterioration of the polymer by activity of the fungi, resulting in a large decrease in resistivity. The data support the hypothesis that polyimides are susceptible to microbial degradation resulting in the corrosion of underlying metal. They also confirm the versatility of EIS as a method in evaluation of the biosusceptibility of polymers. The dielectric properties of polyimides could be altered drastically following growth of microbial biofilms [25, 26, 30, 31, 47], which has wide implications for protection of this class of materials under tropical and subtropical conditions because both humidity and airborne microbial loading are high. This form of deterioration may be slow under ambient conditions in dry and cold regions; however, the deterioration processes can be accelerated in humid conditions or in enclosed environments, for example, submarines, space vehicles, aircraft, and other closed industrial facilities. Very small changes in material properties by the formation of biofilms and trapping of moisture will result in serious functional consequences to the systems. E2.2. Fiber-Reinforced Polymeric Composite Materials. The increasing usage of fiber-reinforced polymeric composite materials (FRPCMs) as structural components of public

428


FIGURE 30.3. Scanning electron micrograph of fungi on fiberreinforced polymeric composite coupon (scale bar, 10 mm).

structures and aerospace applications has generated an urgent need to evaluate the biodegradability of this new class of material. It has become clear that FRPCMs are not immune to the colonization by natural microorganisms, including both fungi and bacteria [4–7, 27–29, 37]. Through a series of investigations, impurities and chemical additives in FRPCMs promote microbial growth as they are sources of carbon and energy (Fig. 30.3, Table 30.1). Research has shown that slowgrowing microorganisms cause much greater damage to materials than the fast-growing microorganisms [11], but no critical assessment has been made in relation to FRPCMs. At least two groups reported microbial degradation of FRPCs [6, 26, 32–35, 149]. Wagner and her collaborators used a mixed culture of microorganisms, including a sulfate-reducing bacterium, commonly used in corrosion tests. In contrast, Gu and colleagues [8, 26–29, 33–35, 71] used a fungal consortium originally enriched from degraded polymers. This consortium consisted of Aspergillus versicolor, Cladosporium cladosorioides, and a Chaetomium sp. Initial physical and mechanical tests were not sufficiently sensitive to detect any significant physical changes in the bulk materials after 120 days exposure [34, 36, 200]. However, physical penetration of fungal hyphae into the resin matrices has been observed with Scanning electron microscopy (SEM) (Fig. 30.4), indicating that resins were being actively degraded, in turn suggesting that the materials were at risk of failure. Natural populations of microorganisms are capable of growing on surfaces of FRPC coupons at both relatively high humidity (65–70%) and lower humidity conditions (55–65%) [6, 28]. The accumulation of bacteria on surfaces of composites develops into a biofilm layer, providing some initial resistance to further environmental changes, for example, drying. Since use of a range of mechanical tests did not detect changes in composite coupons after exposure to a fungal culture [200], EIS was applied to examine changes in

FIGURE 30.4. Scanning electron micrograph of fungi penetrating into matrices of fiber-reinforced polymeric composite coupon (scale bar, 10 mm).

resistivity of composite materials after fungal exposure. EIS indicated a significant decline in material resistivity and an increase in conductance, providing the first demonstration that this technique can be used in monitoring FRPCMs and electrochemical properties can be used to identify microbial attack [33, 35]. The polymeric composites tested included fluorinated polyimide/glass fibers, bismaleimide/graphite fibers, poly(ether-ether-ketone) (PEEK)/graphite fibers, and epoxy/graphite fibers (Table 30.1) [35]. Graphite fibers are very susceptible to biofilm formation and the surface treatment sizing chemicals are utilized by bacteria as a source of carbon and energy (Fig. 30.5). A critical question remains about the effect of FRPCM degradation on the mechanical properties and integrity of these composite materials. It is apparent that EIS is sufficiently sensitive to detect resistivity/conductance of com-

FIGURE 30.5. Scanning electron micrograph of bacteria forming biofilm on surface of graphite fiber (scale bar, 5 mm).

MECHANISMS OF DEGRADATION

posite coupons, but surface-sensitive analytical techniques are needed to register the minor changes correlating to changes in structure and integrity of polymeric matrices. Acoustic techniques have been proposed as a means of detecting changes in the physical properties of the composite [149], but no comparison has been made between EIS and acoustic techniques. E2.3. Protective Coatings. Polymeric coatings are a class of chemicals with increasing production volume and are designed to protect underlying metals from corrosive chemical species and microorganisms. Microbial colonization and formation of biofilms on coatings may accelerate and severely damage the protective coatings and then the underlying metals. Natural bacterial populations were found to readily form microbial biofilms on surfaces of coating materials, including epoxy and polyamide primers and aliphatic polyurethanes (Table 30.1) [6, 28, 36]. Using EIS, both primers and top coatings demonstrated susceptibility to exposure to a mixed culture of fungi compared to sterile controls [28]. Primers were more rapidly degraded in terms of the EIS signal than aliphatic polyurethane coating. A common approach in dealing with coating

429

life is to incorporate biocides in coating formulations, but addition of biocides to polyurethane coatings did not inhibit bacterial attachment or significantly reduce bacterial growth [28, 37]. E2.4. Packaging Polyethylenes. High-density and lowdensity polyethylenes (PEs) are primarily used in product packaging as sheets and thin films. Their degradability in natural environments poses serious environmental problems due to their very slow degradation rate under natural conditions and the hazard they present to freshwater and marine animals. Biodegradation of PEs has been studied extensively [90, 92]. It is believed that polymer additives, such as starch, antioxidants, coloring agents, sensitizers, and plasticizers, may significantly alter the biodegradability of the parent polymers. Degradation rates may be increased by 2–4% following photosensitizer addition. In one study, extracellular culture concentrates of three Streptomyces species were inoculated to starch containing PE films [95, 96], and PE was claimed to be degraded. Degradation of PE by microorganisms may be minimal without taking into account physical and chemical processes, and the data on degradation of PE-containing starch are

TABLE 30.2. Summary of ASTM Methods and Practices and Others for Testing Biodegradation of Polymers and Key Characteristics of Methods ASTM Code

Purposes

D5209-92

Aerobic degradation in municipal sewage sludge Anaerobic degradation in municipal sewage sludge Aerobic biodegradability by specified microorganisms

D5210-92 D5247-92

D5271-92

G21-90

Aerobic biodegradation in activated sludge and wastewater Aerobic biodegradation in composting conditions Resistance to fungi

G22-76

Resistance to bacteria

MIL-STD-810E Method 508.4

Resistance to fungi

D5338-92

Microorganisms Involved and Key Features Indigenous microorganisms in sewage sludgre Indigenous microorganisms in sewage sludgre Streptomyces badius ATCC39117

Streptomyces setonii ATCC39115 Streptomyces viridosporus ATCC 39115 or other organisms agreed upon Municipal sewage treatment plant

2–4-month-old compost Aspergillus niger ATCC 9642 Aureobasidium pullulans ATCC15233 Chaetomium globosum ATCC6205 Gliocladium virens ATCC9645 Penicillum pinophilum ATCC11797 Pseudomonas aeruginosa ATCC 13388 —

Parameters Monitored

References

CO2 evolved

Cited in [46]

CO2 and CH4 evolved

Cited in [46]

Weight loss, tensile strength, elogation and molecular weight distribution

[204]

Oxygen consumption

[205]

Cumulative CO2 production Visual evaluation

[206] [207]

Visual evaluation

[208]

Visual evaluation

[209]

430


questionable. In addition, microbial metabolites may contaminate the PE surfaces and could be misinterpreted as degradation products of the parent PE when chemical characterization is carried out. Abiotic degradation of PE is evidenced by the appearance of carbonyl functional groups in abiotic environments. In contrast, an increase of double bonds was observed when polymers showed weight loss resulting from biodegradation [201]. It was then proposed that microbial PE degradation is a two-step process: an initial abiotic photooxidation followed by a cleavage of the polymer carbon backbone. However, the mechanism of the second step needs extensive analysis before plausible conclusions can be drawn. Lower molecular weight PEs including paraffin can be biodegraded and they undergo hydroxylation oxidatively to form an alcohol group followed by formation of carboxylic acid. At higher temperatures, ketones, alcohols, aldehydes, lactones, and carboxylic acids are formed abiotically within six weeks [201]. PE pipes used in gas distribution systems

may fail due to cracking, but it is unlikely that biological processes are involved [202]. E2.5. Polypropylenes. Polypropylenes (PPs) are widely utilized in engineering pipes and containers. Degradation of PPs results in a decrease of their tensile strength and molecular weight. The degradation mechanism may involve the formation of hydroperoxides, which destabilize the polymeric carbon chain to form a carbonyl group [126, 203]. This step has been confirmed by infrared spectroscopy, but no strong evidence is available for the biochemical basis of degradation of PPs.

F. DEGRADATION METHODOLOGY Traditionally, microbiological degradation of polymeric materials is carried out in microbiological nutrient media for short-term investigations [73, 113, 149]. The American

TABLE 30.3. Comparison of Several Methods Available for Testing Degradability of Different Polymers and Under Range of Environmental and Simulation Conditions Methods

Polymer Forms

Gravimetry

Film or physical intact forms

Respirometry

Inoculum and Degradation Criteria Monitored

Comments

References

A wide range of inocula can be used, from soil, waters, sewage, or pure species of microorganisms from culture collections.

This method is robust and also good for isolation of degradative microorganisms from environments of interest. Reproducibility is high. Disintegration of polymer cannot be differentiated from biodegradation.

[40–46, 83, 84, 92, 93]

Film, powder, liquid, and virtually all forms and shapes

Either oxygen consumed or CO2 produced under aerobic conditions. Under methanogenic conditions, produced methane can be monitored.

This method is most adaptable to a wide range of materials. It may require specialized instruments. When fermentation is the major mechanism of degradation, this method gives an underestimation of the results.

[40–46, 204–209]

Surface hydrolysis

Films or other

Generally aerobic conditions, pure enzymes are used. Hydrogen ions (pH) released are monitored as incubation progresses.

Prior information about the degradation of the polymer by microorganims or particular enzymes is needed to target the specific test.

[180, 181]

Electrochemical impedance spectroscopy

Films or coatings resistant to water

The test polymers should be adhered on the surface of conductive materials and electrochemical conductance is recorded.

Polymer must initially be water impermeable for signal transduction. Degradation can proceed quickly, and as soon as degradation is registered, no further degradation processes can be distinguished.

[26, 28, 30–35]

USE OF BIOCIDES

Society for Testing and Materials (ASTM) offers a range of standard methodologies for assessment of material degradation under specific conditions [6, 46, 204–208] (Table 30.2). These methods are widely accepted by chemists and engineers but few biologists. However, the guidelines (1) are designed for a specific set of conditions which can be reproduced anywhere but may not be relevant to an individual environment; (2) fail to take into account microbiological processes and subsequent biochemical characterization; and (3) are usually qualitative with no or minimal emphasis on the mechanisms involved during degradation (Table 30.3). Materials used for aerospace applications have been assessed in the United States by application of MIL-STD810E method 508.4 [209]. In addition to the problems described above, a weakness of this method includes the failure to determine the effects of long-term exposure of the materials, as slow-growing microorganisms appear to be more effective in degrading polymeric materials than fast growers [210]. This method also emphasizes fungi and ignores any participation of bacteria in degradation processes even though both groups of microorganisms may play synergistic roles in the complete degradation of complex polymeric materials.

G. PLASTICIZERS Polymeric materials are known to contain a wide array of chemicals, including plasticizers to improve their processTABLE 30.4.

431

ibility and product quality; these plasticizers mostly include phthalate esters, specifically dimethyl phthalate esters (namely ortho-dimethyl phthalate ester, dimethyl isophthalate ester, and dimethyl terephthalate ester), di-n-butyl phthalate ester, and dibutylbenzyl phthalate ester [12, 101–112]. Since these chemicals are not covalently bound to the polymer resins, they can be utilized by many bacteria isolated from activated sludge [211], mangrove sediments [101–105, 108] and deep-ocean sediment [12, 106, 107] (Table 30.4). Biochemical cooperation is required between two different bacteria, namely Arthrobacter species and Sphingomonas paucimobilis in utilization of orthodimethyl phthalate ester (Fig. 30.6) [105, 107], but Variovorax paradoxus T4 isolated from deep-ocean sediment of the South China Sea is capable of utilizing dimethyl terephthalate ester as sole carbon and energy source (Fig. 30.7). Di-n-butyl phthalate dibutylbenzylphthalate esters can be degraded by Pseudomonas fluorescence B-1 isolated from mangrove sediment [109–112]. Results collectively indicated that enzymes involved in the initial cleavage of the two ester bonds can be highly selective.

H. USE OF BIOCIDES Microorganisms are commonly controlled by application of biocides. While most antimicrobial products are effective against the growth of microorganisms in liquids, biofilm bacteria are more resistant and can rapidly develop resistance after exposure to chemicals [28, 37, 128,

Microorganisms and Degradation of Selective Plasticizers in Class of Phthalate Diester with Source of Bacteria

Substrate

Microorganism(s) Involved

Source of Inoculum

Degradation Intermediates

References

Dimethyl phthalate

Comamonas acidovorans fy-1, Xanthomonas maltophila, and Sphingomonas paucimobilis Rhodococcus ruber Sa

Activated sludge

Monomethyl phthalate, phthalic acid

[221, 222]

Mangrove sediment

[101, 102]

Fusarium sp.

Mangrove sediment

Klebsiella oxytoca Sc and Methylobacterium mesophilicum Sr Rhodococcus ruber Sa

Mangrove sediment

Variovorax paradoxus Rhodococcus ruber Sa

Deep-ocean sediment Mangrove sediment

Pasteurella multocida Sa

Mangrove sediment

Di-n-butyl phthalate

Variovorax paradoxus Sphingomonas yanoikuyae Pseudomonas fluorescens B-1

Deep-ocean sediment Deep-ocean sediment Mangrove sediment

butylbenzylphthalate

Pseudomonas fluorescens B-1

Mangrove sediment

Monomethyl phthalate, phthalic acid Monomethyl phthalate phthalate Monoisophthalate, phthalic acid Monoisophthalate, phthalic acid Monoisophthalate Monoterephthalate, phthalic acid Monoterephthalate, phthalic acid Monoterephthalate, ? Monobutyl phthalate, phthalic acid benzylphthalate, phthalic acid

Dimethyl isophthalate

Dimethyl terephthalate

Mangrove sediment

[220] [103–105] [12,106, 107] [101] [101, 102] [103] [12, 106, 107]. [107] [109–111] [109–111]

432


O

O

O

OH

OH

OCH3

OH

S. paucimobilis

O

OCH3

O

Arthrobacter sp.

OCH3

Further Degradation

Arthrobacter sp. or S. paucimobilis

O

O

OH Further Degradation

OH

Arthrobacter sp. or S. paucimobilis

O

FIGURE 30.6. Biochemical pathway for degradation of ortho-phthalate diester by Arthrobacter species and Sphingomonas paucimobilis through metabolic collaboration.

212–215]. Biofilms reduce the effectiveness of the biocides because of reduced diffusion rates and the binding affinity of exopolymeric materials in the biofilm for biocides. In addition, microorganisms have plasmids which can be effectively exchanged between organisms to rapidly build resistance to the environmental chemicals and biocides [213]. Most industrial tests are conducted in the liquid phase, which does not represent real environmental conditions of exposure. For example, a biocide that effectively inhibits bacterial growth in a test involving liquid culture failed to control growth on surfaces of a polyurethane coating [4, 6, 28, 37]. It is strongly recommended that tests for biocide efficacy should be applied to material surfaces. In addition, compatibility of a biocide with a material should be considered.

(a)

O

OCH 3

O

It is routine practice that corrosion protective paint formulations are amended with biocides to prolong the shelf-life of products and extend service time [216]. Industrial development of water-based paints and coatings creates the problem of increased microbial contamination and degradation (Fig. 30.8), because these environmentally friendly chemicals may be susceptible to rapid colonization and utilization by microorganisms. Microbial growth deteriorates the quality of products and also provides initial contamination to the product in service. The basic approach to such a problem should emphasize environmental control so that microorganisms are prevented from attaching to surfaces in the first place. Conditions that promote microbial growth should also be avoided so that microbial population growth is kept to a minimum.

OH

O

OH

CO 2 + H 2 O

3 HCO

3 HCO

O

dimethyl terephthalate (DMT)

(b)

O

3 HCO

OCH

3

O

dimethyl terephthalate (DMT)

O

monomethyl terephthalate (MMT) O

3 HCO

OH

O

monomethyl terephthalate (MMT)

HO

O

terephthalate acid (TA) O

HO

OH

O

terephthalate acid (TA)

FIGURE 30.7. Biochemical degradation pathways of dimethyl terephthlate by (a) Variovorax paradoxus T4 and (b) Sphingomonas yanoikuyae DOS01 isolated from deep-ocean sediment.

REFERENCES

433

ACKNOWLEDGMENTS We thank Zhenye Zhao for drawing Figure 30.1; Ed Seling at the Museum of Comparative Zoology, Harvard University, for the scanning electron micrographs taken and used here in Figures 30.2–30.5; and Jessie Lai for assistance in organizing the final figures and archiving. REFERENCES

FIGURE 30.8. Photograph of two flasks containing water-based aliphatic polyurethane coating inoculated with soil as inoculums (right) and kept sterile throughout incubation period (left).

I. BIODEGRADATION AND HEALTH Polymeric materials are increasingly used not only in engineering and structures, building, and decoration but also in medical implants. Material types cover a wide range of composition, including absorbable sutures and artificial skin. The materials must be thoroughly examined and tested for potential biofilm formation. Some additives, such as PEGs, can promote microbial growth by providing readily available sources of carbon and energy. Other chemicals may serve as attractants for bacteria. In addition, microbial metabolites during active growth, particularly secondary ones, can be allergens or irritants to humans in building materials [217, 218]. Fungi have been studied for their tolerance to lower humidity, where bacteria are barely capable of survival or growth [28]. The growth of fungi may not only damage the materials, integrity but also produce toxic secondary metabolites [8, 219].

J. CONCLUSIONS Polymeric materials are a diverse class of polymers with multiple applications. Assessment of their biodegradability must address the relationship between their chemical structures and proposed applications. An array of test protocols are available and should be developed to detect early changes in polymer physical and chemical properties. EIS is a sensitive method for detecting early physical changes in polymer films, but the microbial role in deterioration needs to be confirmed with more sensitive methods. In most cases, accelerated testing under specific environmental conditions will be required to determine susceptibility to microbial deterioration and to assess protective measures.

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31 DURABILITY OF CONCRETE* V. M. MALHOTRA Consultant, Ottawa, Ontario, Canada

A. Introduction B. Deterioration caused by freezing and thawing cycles B1. Mechanisms of freezing and thawing deterioration B1.1. Hydraulic pressure B1.2. Osmotic pressure B1.3. Measures to avoid damage due to freezing and thawing cycling B2. Mechanisms of surface scaling due to combined action of freezing and thawing cycling and application of deicing salts B2.1. Layer-by-layer freezing B2.2. Thermal shock B2.3. Supercooling due to preventive salt application B2.4. Osmotic pressure B2.5. Measures to avoid damage due to deicing salt scaling C. Deterioration caused by alkali–aggregate reactions C1. Types of reaction C2. Nature of reaction and expansion processes C3. Conditions conducive to alkali–aggregate reactivity C3.1. Inherent reactivity of siliceous material C3.2. Total alkali content of concrete C3.3. Environmental considerations C3.4. Preventive measures against alkali–aggregate reactivity C4. Use of nonreactive aggregates C5. Use of low-alkali cement C6. Limiting alkali content of concrete mixtures C7. Use of supplementary cementing materials D. Deterioration due to sulfate attack D1. Mechanisms of sulfate attack D2. Measures to protect against sulfate attack D2.1. General measures Her Majesty the Queen in Right of Canada, as represented by the Minister of Natural Resources, 2010.

E. Deterioration of concrete in seawater E1. Requirements of concrete for marine exposure F. Carbonation of concrete F1. Mechanism of carbonation F2. Factors affecting carbonation F3. Rate of carbonation F4. Carbonation of concrete-containing supplementary cementing materials F5. Measures to prevent carbonation G. Concluding remarks References Bibliography

A. INTRODUCTION According to the American Concrete Institute (ACI) Committee 201, durability of Portland cement concrete is defined as its ability to resist weathering action, chemical attack, abrasion, or any other process of deterioration; that is, durable concrete will retain its original form, quality, and serviceability when exposed to its environment. If concrete is proportioned correctly and is cast, consolidated, and cured properly, it should be maintenance free for a very long time. But in practice this rarely happens with the result that concrete structures deteriorate prematurely and need constant repairs and maintenance at a considerable cost. This chapter discusses selected aspects of physical and chemical attacks that affect significantly the long-term durability of concrete. The physical attack that causes major distress in concrete structures is the freezing and thawing phenomenon; the chemical attacks include distress caused


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DURABILITY OF CONCRETE

by the sulfates, alkali-aggregate reactions, seawater, and carbonation of concrete. Before discussing in some detail the deterioration of concrete by physical and chemical mechanisms, it should be stressed that porosity of hydrated cement paste, and hence concrete, plays a very major role in its durability. As porosity of concrete is directly related to its water–cement (W/C) ratio, any decrease in W/C will reduce significantly the porosity and hence its permeability. The decreased permeability will decrease the transportation of aggressive chemicals into concrete and will also control the moisture content during environmental changes. For example, a decrease in W/C concrete from 0.80 to 0.40 will reduce the coefficient of water permeability from 130 10 12 to 10 10 l2 m/s.

when the bulk of the concrete is affected. The concrete made with sound aggregates can be fully protected against freezing and thawing by proper air entrainment; however, air entrainment is found rarely to protect concrete totally from damage due to the combined action of freezing and thawing cycling and the deicing concrete chemicals. It is generally agreed upon that the deterioration caused by the deicing chemicals is mostly physical in nature and that the chemical reactions of the salts with the cement hydration products play only a secondary role in the deterioration mechanisms. Figure 31.1 shows a concrete structure damaged by freezing and thawing cycling.

B. DETERIORATION CAUSED BY FREEZING AND THAWING CYCLES

B1.1. Hydraulic Pressure. Water in the capillary pores of cement paste in concrete expands 9% upon freezing. If the increased volume is smaller than the space available, no damage will occur; otherwise the excess water will be expelled by the hydraulic pressure. As cement paste is a permeable material, there is a possibility that excess water can escape from the capillary pores to the nearest air void during the process of freezing if the air void is unfilled. An entrained-air void will be unfilled, that is, filled with air unless a crack has penetrated it. Water forced into it by the mechanism described will exit immediately on thawing, forced out by the compression of the air, and hence no liquid will remain to evaporate and form secondary deposits. The presence of such deposits indicates that the bubble has been penetrated by a crack. Thus, the magnitude of this hydraulic

In Canada, northern parts of the United States, northern Europe, Japan, Korea, Russia, and northern parts of China, concrete is subjected to repeated cycles of freezing and thawing. In addition, in several of these countries, deicing chemicals, such as sodium chloride, are routinely used to melt ice and snow on the highways, roads, and sidewalks. The freezing and thawing cycling and the combined action of freezing and thawing cycling and the application of the deicing salts result in considerable deterioration of concrete. In the deicing salt scaling of concrete, the mortar near the surface flakes or peels away. It is primarily a surface phenomenon, in contrast to the internal cracking of concrete in freezing and thawing cycling

B1. Mechanisms of Freezing and Thawing Deterioration

FIGURE 31.1. Damage of a concrete structure caused by freezing and thawing cycling [6].

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441

pressure depends on the permeability of the cement paste, the distance that water must travel to reach the nearest unfilled void, the rate of freezing, the rate of ice formation, and the degree of saturation of the paste. If the pressure is high enough to stress the surrounding paste beyond its tensile strength, it will cause cracking [1, 2].

230 mm with no detrimental effects on the performance of concrete subjected to freezing and thawing cycling.

B1.2. Osmotic Pressure. In addition to hydraulic pressure caused by water freezing in capillary cavities, the osmotic pressure resulting from partial freezing of the solutions in such cavities can be another source of destructive expansions in cement paste. Water in the capillary cavities is not pure; it contains various soluble substances. Such solutions freeze at slightly lower temperatures than pure water. When solutions of different concentrations of soluble materials are separated by a permeable barrier, the solvent particles move through the barrier toward the solution of greater concentration. The existence of local salt concentration gradients between capillaries is envisaged as the source of osmotic pressure [3]. The hydraulic pressure due to an increase in the specific volume of water on freezing in large cavities and the osmotic pressure due to salt concentration differences in the pore fluid do not appear to be the only causes of expansion of cement paste exposed to frost action [4], but they are believed to be the main ones.

In addition to the hydraulic and osmotic pressure theories discussed above, there are a few more phenomena that are associated with the presence of deicing chemicals and are believed to contribute to the surface-scaling type of deterioration of concrete (Fig. 31.2).

B1.3. Measures to Avoid Damage Due to Freezing and Thawing Cycling. Air entrainment in concrete has been used as a means of reducing internal damage due to frost action. In air-entrained Portland cement paste, every air void is assumed to be bordered by a zone in which the hydraulic pressure cannot become high enough to cause damage. By reducing the distance between the voids to the point where the protected zones overlap, the generation of disruptive hydraulic pressures during the freezing of water in the capillary cavities can be prevented. In order to avoid damage due to freezing and thawing cycling, concrete in North America is routinely air entrained using chemical admixtures. At a given air content, the protection provided by the air voids against repeated cycles of freezing and thawing is usually greater, the larger the number of voids per unit volume of paste. This implies that voids are more effective when they are close together. It is generally agreed upon that the cement paste is well protected against of the airthe effects of frost action if the spacing factor L void system is 200 mm or less, as determined in accordance with American Society for Testing and Materials (ASTM) C 457 test procedure. Field experience has shown that concrete incorporating between 5 and 9% air by volume which are of the order of will generally yield values of L 200 mm. When supplementary cementing materials such as fly ash or blast-furnace slag are incorporated into concrete can reach together with superplasticizers, the value of L

B2. Mechanisms of Surface Scaling Due to Combined Action of Freezing and Thawing Cycling and Application of Deicing Salts

B2.1. Layer-by-Layer Freezing. Differences in the salt concentration in concrete lead to corresponding differences in freezing temperatures of various layers. When ice formation occurs in such a layer-by-layer way, stresses can develop whose extent depends on the dilation difference between the frozen and unfrozen layers [5]. B2.2. Thermal Shock. Heat is required for the melting of ice and snow. When thawing takes place by the application of deicing chemicals, the heat is extracted mostly from the concrete. The great heat loss causes a rapid temperature drop on the surface of the concrete. The temperature gradient developed can cause stresses of short duration near the surface that may exceed the tensile strength of the concrete and cause damage [5]. B2.3. Supercooling Due to Preventive Salt Application. The use of deicing chemicals prevents water from freezing

FIGURE 31.2. Surface scaling of concrete caused by deicing chemicals [6].

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on the surface of concrete at a temperature near 0 C. When the supercooled water near the surface of the concrete eventually freezes, the destructive effect of the phase transition will be greater than with normal freezing [5].

deleterious expansion and consequent cracking of the concrete or mortar may result [7].

B2.4. Osmotic Pressure. The use of deicing chemicals will increase the concentration of the chemicals in the capillary pores near the surface of the concrete, which may build up an osmotic pressure high enough to cause a rupture of the cement paste near the surface.

The reaction starts when the alkaline hydroxides in the concrete fluid (Na–OH, K–OH) attack the surface of silica minerals in the aggregate. This results in an alkali silica gel and an alteration of the aggregate surface, known as “reaction rim.” The resultant gel has a strong affinity for water and, consequently, a tendency to increase in volume. The expanding gel exerts an internal pressure that is sufficient to fracture the surrounding cement paste. Some of the softer gel can leach out through the voids and cracks, but expansion of the solid and semisolid products are more damaging. Figure 31.3 shows a damaged bridge structure due to the alkali–silica reaction, and Figure 31.4 shows expansion in a concrete median barrier due to the alkali– aggregate reactivity. Table 31.1 gives a list of minerals and of rock types that have been found reactive in concrete structures. The information given in this table should, however, be used with care. For example, even if some greywackes and granites have been found to be reactive in certain parts of Canada, it does not mean that all greywackes and granites found in Canada are reactive or will react to the same degree. Petrographic analysis will generally help to identify and to determine the

B2.5. Measures to Avoid Damage Due to Deicing Salt Scaling. So far, there is no agreement that any of the above mechanisms are the primary reasons for the deterioration of concrete due to scaling. Consequently, no effective solution has been possible to prevent or to reduce the surface scaling of concrete. Research and practical experience indicate that air entrainment in concrete is effective in preventing internal deterioration; however, the scaling is never completely prevented by the air entrainment, and the critical air-void concept may not be applicable to deterispacing factor (L) oration due to the scaling. According to Cordon [6], field experience has shown that linseed oil acts as an antiscaling agent. The solutions of this oil and flammable solvents are sprayed on highways, streets, and bridges to reduce damage caused by deicing salts.

C2. Nature of Reaction and Expansion Processes

C. DETERIORATION CAUSED BY ALKALI–AGGREGATE REACTIONS For many years, aggregates were believed to be essentially inert and chemically nonreactive in concrete mixtures. Unfortunately, this often is not true, and deleterious reactive aggregates have been found in many parts of the world, including Africa, Australia, Canada, China, England, India, Japan, New Zealand, the Scandinavian countries, and the United States. C1. Types of Reaction There are two types of alkali–aggregate reactions:

FIGURE 31.3. Cracking in a concrete bridge caused by alkali–silica reaction [8].

(a) Alkali–silica reaction (b) Alkali–carbonate reaction Of these two reactions, the alkali–silica reaction is the most common. For this reason, the alkali–carbonate reaction will not be discussed here. The alkali–silica reaction is a reaction in either mortar or concrete between the hydroxyl (OH ) ions associated with the alkalies (Na2O and K2O) from the cement or other sources and certain mineral phases that may be present in the coarse or fine aggregates; under certain conditions,

FIGURE 31.4. Cracking of a highway medium barrier caused by alkali–aggregate reactivity [8].

DETERIORATION CAUSED BY ALKALI–AGGREGATE REACTIONS

TABLE 31.1. Mineral Phases Susceptible to Deleterious Reactions with Alkalies in Cement and Corresponding Rocksa Alkali-Reactive Silica Minerals and Rocks B1.1 Alkali-Reactive Silica Minerals and Volcanic Glasses (Classical Alkali–Silica Reaction) Reactants Rocks

Opal, tridymite, cristobalite; acid, intermediate, and basic volcanic glasses; artificial glasses; beekite Rock-types containing opal such as shales, sandstones, silicified carbonate rocks, some cherts and flints, diatomites Vitrophyric volcanic rocks; acid, intermediate and basic, such as rhyolites, dacites, latites, andesites, and their tuffs; perlites, obsidians; all varieties with a glassy goundmass; some basalts B1.2 Alkali-Reactive Quartz-Bearing Rocks

Reactants

Rocks

a

Chalcedony, cryptocrystalline to microcrystalline and macrogranular quartz with deformed crystal lattice, rich in inclusions, intensively fractured or granulated; poorly crystalline quartz at grain boundaries Cherts, flints, vein quartz, quartzite, quartz-arenite, quartzitic sandstones that contain microcrystalline to cryptocrystalline quartz, and /or chalcedony Volcanic rocks such as in B1.1 (above) but with devitrified, crypto- to microcrystalline groundmass Micro- to macrogranular silicate rocks of various origin: Metamorphic rocks: gneisses, quartz-mica schists, quartzites, homfilses, phyllites, argillites, slates Igneous rocks: granites, granodiorites, charnockites Sedimentary rocks: sandstones, greywackes, siltstones, shales, siliceous limestones, arenites, arkoses Sedimentary rocks (sandstones) with epitaxic quartz cement overgrowth

From Canadian Standards Association document CSA 23.1 Appendix B.

proportion of the various potentially reactive rock types in an aggregate sample from a quarried operation or a gravel deposit. However, this information alone will not permit the prediction of the magnitude of the reaction that may occur in the field. C3. Conditions Conducive to Alkali–Aggregate Reactivity The rate and extent of expansion due to alkali–aggregate reaction are affected by a large number of factors. These can generally be grouped as follows [8]: (a) Inherent reactivity of the siliceous material (b) Total alkali content of concrete (c) Environmental considerations

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C3.1. Inherent Reactivity of Siliceous Material. The nature, amount, and particle size of the siliceous phase within the rock particles play a major role on the actual rate of alkali–aggregate reaction in mortar or concrete. Poorly crystalline forms of silica such as opal and volcanic glass are very reactive because of the nature of their structure through which the alkali hydroxide ions (Na–OH and K–OH) can penetrate quickly. As little as 1% of such reactive components in an aggregate may lead to deleterious reaction and cracking and generally within 10 years of construction. Quartz is one of the major constituents of the various rocks found in the earth’s crust. Quartz has a well-organized threedimensional crystalline structure that reacts at a much slower rate than opal and volcanic glasses. However, very small particles of quartz, because of their increased surface area, or larger particles, which have been subjected to stress in their geological history and consequently show defects in their crystalline structure, may be fairly reactive. C3.2. Total Alkali Content of Concrete. The source of alkali for the alkali–aggregate reaction is generally considered to be derived from the Portland cement in concrete. Alkalies may also be contributed by aggregates, admixtures, supplementary cementing materials, and extraneous sources such as deicing salts and seawater. The amount of alkalies contributed by the cement for the alkali–aggregate reaction is calculated by adding the sodium oxide content to 0.658 times the potassium oxide content. The calculation provides the sodium oxide equivalent. The minimum alkali content at which the alkali–aggregate reaction will occur will vary from one aggregate to another and the conditions to which the concrete is subjected. The alkali–aggregate reaction will generally not occur when the total alkali content is B MF (ends exposed) > hot-pressed (HP) SiC > B MF (circumferential surface exposed) > Si. It should also be noted that ceramic reinforcements may vary in purity and structure, and some reinforcements are in themselves composites. For example, SiC MFs have carbon-rich outer layers and carbon cores, and their polarization diagrams have a stronger

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FIGURE 35.4. Collection of anodic polarization diagram of ultrapure Al (99.999%) [32], pure Al (99.9%) [33], Al–2 wt % Cu [34], Al 6061-T6 [32], ultrapure Ti (99.99%) [35, 36], Ti 15-3 [35, 36], ultrapure Cu (99.999%), pure Mg (99.95%) [37], and Mg ZE 41A [37]. Cathodic polarization diagrams of P100 GrE [32], HP SiC [32], SiCMFE [38], SiCMFS [38], Si [33], TiB2 [32], BMFE [31], and BMFS [31] exposed to aerated 3.15 wt % NaCl at 30 C.

resemblance to P100 Gr than HP SiC. In addition, since the MFs often have cores that are of a different material than their surfaces, the orientation of the reinforcements may also affect corrosion behavior. The polarization behavior of MF electrodes with the circumferential surface exposed (MFS) is different compared to the behavior of MF electrodes with the ends exposed baring the cores (MFE). Compare cathodic curves for SiC MFS versus SiC MFE and B MFS versus B MFE in aerated (Fig. 35.4) and deaerated (Fig. 35.5) 3.15 wt % NaCl. C1.4. Reinforcement Photoelectrochemistry. If the MMC reinforcements or constituents are semiconductors, galvanic currents between the matrix metal and the semiconductor could be suppressed under illumination if the semiconductor is n type or accelerated if the semiconductor is p type. An n-type semiconductor is photoanodic and can promote photooxidation reactions under illumination. One such reaction is the oxidation of water. Hence, when an MMC containing n-type semiconductors is wet and under illumination, photogenerated electrons from the n-type semiconductor could polarize the MMC to more negative potentials inducing cathodic protection [39]. Accordingly, anodic

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were immersed in air-exposed 0.5 M Na2SO4 solutions, cathodic current densities and open-circuit potentials increased sharply during illumination. In these MMCs exposed to the outdoors, corrosion films were also thicker on the sunlit surfaces as apposed to the shaded surfaces [41]. C1.5. Reinforcement Area Fraction. If the galvaniccorrosion rate is under cathodic control, galvanic corrosion should increase as the reinforcement area fraction increases. The catchment area principle [42] can be used to determine iGALV as a function of the area fraction of the cathodic reinforcement [32]: iGALV ¼ iC

FIGURE 35.5. Collection of anodic polarization diagram of ultrapure Al (99.999%) [32], pure Al (99.9%) [33], Al–2 wt % Cu [34], Al 6061-T6 [32], ultrapure Ti (99.99%) [35, 36], Ti 15-3 [35, 36], ultrapure Cu (99.999%), pure Mg (99.95%) [37], and Mg ZE 41A [37]. Cathodic polarization diagrams of P100 GrE [32], HP SiC [32], SiCMFE [38], SiCMFS [38], Si [33], TiB2 [32], BMFE [31], and BMFS [31] exposed to deaerated 3.15 wt % NaCl at 30 C.

current densities on Al 6092/Al2O3/20 P-T6 MMCs immersed in air-exposed 0.5 M Na2SO4 solutions increased sharply under illumination, which was attributed to photoanodic currents generated by water oxidation on n-type TiO2 particles in the microstructure that were likely introduced with the Al2O3 reinforcements [40]. The open-circuit potentials also decreased upon illumination, indicating that the n-type TiO2 induced cathodic protection. Outdoor exposures of these Al 6092/Al2O3/20 P-T6 MMCs showed that corrosion films were thinner on the topside of specimens exposed to sunlight as compared to the backside of the specimens not exposed to sunlight [41]. Interestingly, MMCs containing p-type semiconductors had thicker corrosion films on the sunlit surfaces as apposed to the shaded surfaces [41]. A p-type semiconductor is photocathodic and under illumination promotes photoreduction reactions. Depending on the electrolyte conditions, proton or oxygen reduction may be enhanced at the p-type semiconductor. In the presence of moisture and illumination on MMCs that contain p-type semiconductors, photoreduction causes the cathodic current to increase, raising the corrosion potential and inducing greater dissolution of the matrix [39]. Accordingly, during cathodic polarization of Al 6092/SiC/P-T6 MMCs [39] that

XC 1 XC

ð35:1Þ

where the parameter iGALV is the dissolution current density of the matrix (i.e., IGALV/anode area), iC is the current density of the cathode, XC is the area fraction of the cathode, and 1 XC is the area fraction of the anode. The value of iC can be set equal to the current density of the cathodic constituents at the galvanic couple potential. For Cu (99.999%), Al (99.999%, 99.9%, Al–2 wt % Cu, and 6061-T6 Al), and Mg (99.95%, ZE41A) coupled to P100 Gr, the iC values are approximately 10 5, 2.5 10 4, and 7.9 10 4, respectively, in aerated 3.15 wt % NaCl at 30 C (Fig. 35.4). By plotting Eq. 35.1, a graph (Fig. 35.6) was generated from which iGALV of Cu, Al, and Mg can be obtained as a function of the area fraction of P100 GrE for exposure to aerated 3.15 wt % NaCl at 30 C. Note that iC for Cu was in the Tafel regime and that for Al and Mg was in the diffusion-limited oxygen reduction regime (Fig. 35.4). Hence, galvanic corrosion rates for Al and Mg could increase significantly with convection, whereas that for Cu should not. The diagram

FIGURE 35.6. Graphs showing galvanic corrosion current density igalv of Mg [i.e., pure Mg (99.95 %), Mg AZ41A], Al [i.e., ultrapure Al (99.999%), pure Al (99.9%), Al–2 wt %Cu-T6, or Al 6061-T6], and ultrapure Cu (99.999%) as a function of the area fraction of XC of P100 GrE in aerated 3.15 wt % NaCl at 30 C. Normal corrosion rates of pure Mg, ultrapure Al, and ultrapure Cu are also plotted as horizontal lines.

FACTORS INFLUENCING MMC CORROSION

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(Fig. 35.6) also shows the normal corrosion rates of ultrapure Cu, ultrapure Al, and pure Mg. Hence, the area fraction of P100 Gr at which the galvanic corrosion rate would exceed the normal corrosion rate of the matrix metal varies with the type of metal. C1.6. Matrix Metal. The matrix alloy plays a significant role in galvanic corrosion. In matrix alloys with noble opencircuit potentials, the effect of galvanic corrosion is generally attenuated. Notice that in deaerated 3.15 wt % NaCl (Fig. 35.5), titanium and copper are virtually immune to galvanic corrosion, whereas magnesium is highly susceptible. C1.7. Environment. The environment can also have a significant effect on galvanic corrosion rates. For matrix alloys that corrode galvanically under cathodic control, galvanic corrosion rates can increase markedly when dissolved oxygen levels increase. The anodic polarization diagrams of several alloys are plotted with the cathodic polarization diagrams of various reinforcements in aerated (Fig. 35.4) and deaerated (Fig. 35.5) 3.15 wt % NaCl at 30 C. Marked increases in galvanic corrosion rates result from aeration (compare Figs. 35.4 and 35.5) for the various types of aluminum when they are galvanically coupled to P100 graphite, SiC MF, B MF, HP SiC, and TiB2. Notice that galvanic corrosion rates do not increase for the aluminums coupled to Si since the polarization curves intersect in the passive aluminum regime, making galvanic corrosion under anodic control. In halide-free environments such as in 0.5 M Na2SO4, aluminum passivates and galvanic corrosion is also under anodic control; hence, aeration (Fig. 35.7) would not increase galvanic corrosion rates higher than that of the passive aluminum current density based on the polarization diagrams. The breakdown of passivity due to the presence of aggressive ions such as chlorides can also significantly increase galvanic corrosion rates (compare Figs. 35.4 and 35.7). C1.8. Microstructure. The corrosion of the MMC is also affected by the physical presence of the reinforcements, even if they are inert. As the matrix alloy corrodes, the reinforcements are often left in relief, leaving behind a network of fissures that trap corrosion products and exacerbate corrosion. The initiation and propagation of corrosion sites are generally influenced by the electrical resistivity and volume fraction of the MMC constituents, including the reinforcements, interphases, and intermetallics. The corrosion behavior of MMCs in the open-circuit condition can be quite different from what might be expected based on their anodic polarization diagrams. For example, in near-neutral 0.5 M Na2SO4 solutions, various aluminum MMCs passivate during anodic polarization [43, 44], but in the open-circuit condition, the same MMCs are susceptible to localized

FIGURE 35.7. Collection of anodic polarization diagram of ultrapure Al (99.999%) [32], pure Al (99.9%) [33], Al–2 wt % Cu [34], Al 6061-T6 [32], pure Mg (99.95%) [37], and Mg ZE 41A [37]. Cathodic polarization diagrams of P100 GrE [46], HP SiC [46], SiCMFE (oxygenated solution) [47], SiCMFS (oxygenated solution) [47], Si [33], and TiB2 [46] exposed to aerated 0.5 M Na2SO4 at 30 C.

corrosion. At the open-circuit potential, localized cathodic sites can become alkaline, and localized anodic sites can become acidified [45]. If the matrix metal is amphoteric like that of aluminum, dissolution may occur in the alkaline cathodic and acidic anodic regions [45]. Corrosion at alkaline cathodic regions usually occurs in aluminum MMCs when the cathodic constituents are of low area fraction, which concentrates the cathodic current and hydroxide ion buildup over a few cathodic sites. Hence, when predicting MMC corrosion behavior, polarization studies alone may be misleading. An example of where corrosion was induced by localized cathodic currents was in Al 6092/Al2O3/20 P-T6 MMCs. In these MMCs, the Al2O3 reinforcement particles are insulators and cannot serve as cathodes; however, the microstructure also contained several types of titanium suboxides with compositions close to that of Ti6O, Ti3O, Ti2O, and TiO; TiO2 (likely doped); Ti–Zr–Al-containing oxides; and Fe–Si–Al intermetallics [40]. Of these particles, the titanium suboxides, TiO2, and the Fe–Si–Al intermetallics supported significant cathodic activity [40]. The area fraction of the non-Al2O3 particles in the MMCs was estimated to be on the order of 0.01 using image analysis [43]. In the open-circuit condition, these sites were observed to be corrosion initiation sites in aerated 0.5 M Na2SO4, where the MMC passivates during anodic polarization.

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C2. Electrochemical Effects of Interphases

C4. Secondary Effects

During the fabrication processing of MMCs, reactions between the reinforcement and matrix may lead to the formation of an interphase at the reinforcement–matrix interface. The presence of the interphase may lead to corrosion behavior different from what might be expected based on virgin MMC constituents. For example, Pohlman [48] could not measure galvanic currents between virgin B MFs and Al 2024 or Al 6061 in 3.5% NaCl solutions, indicating that galvanic corrosion between aluminum matrices and B MFs should be negligible. In actual Al/B/MF MMCs, however, galvanic corrosion takes place between the aluminum matrix and the aluminum boride interphase on the surface layers of the B MFs [48]. Pohlman measured galvanic currents between the aluminum alloys and B MFs that were extracted from the matrix. A 4-mm-thick layer of aluminum boride enveloped the extracted B MFs. Galvanic currents measured between the aluminum alloys and aluminum boride were similar to those between the alloys and the extracted B MFs. When the layer of aluminum boride was removed from the extracted B MFs, the galvanic current ceased, which indicated that the aluminum boride interphase was necessary for galvanic corrosion.

Corrosion in MMCs is also influenced by factors that are not directly caused by the reinforcement, matrix, or interphase but result due to their presence. For example, the reinforcement phase in the MMC may alter the microstructural features in the matrix metal in ways that do not occur in the monolithic matrix alloy. Examples discussed below are the effects of intermetallic phases which may form around the reinforcements by solute rejection during solidification [56], dislocation generation by the mismatch in CTE between the reinforcement and matrix [57], and the susceptibility of the MMC microstructure to physical damage. In addition, problems related to processing deficiencies are disused.

C3. Chemical Degradation of Reinforcements and Interphases MMCs may also degrade by chemical reactions that cannot be directly assessed by electrochemical methods. Reinforcement phases and interphases may undergo chemical degradation which cannot be detected by polarization techniques. Muscovite mica KAl3Si3O10(OH)2 particles of approximately 70 mm in size [49] used in Al MMCs appeared to have absorbed moisture, swelled, and then exfoliated [50] during exposure to non-deaerated 3.5 wt % NaCl solutions. The aluminum carbide Al4C3 interphase which is present in some aluminum MMCs may hydrolyze in the presence of moisture, forming methane and aluminum hydroxide. The rate of Al4C3 hydrolysis was measured to be approximately 1% per hour for hot-pressed Al4C3 (78% of theoretical density and porous) exposed to pure water at 30 C [46]. Methane evolution has been detected from Al/Gr MMCs containing Al4C3 [51, 52]. Buonanno [52] reported that Al4C3 hydrolysis in Al/Gr MMCs leaves fissures at fiber–matrix interfaces. The hydrolysis of Al4C3 therefore could result in rapid penetration into the MMC microstructure through reinforcement–matrix interfaces, leading to the formation of microcrevices. Al4C3 can form by the reaction of aluminum and carbon [53], SiC [54], or B4C [55] at elevated temperatures that could be encountered during MMC processing. Hence, if processing conditions are not properly controlled, MMCs containing Al and carbon compounds could contain this deleterious interphase.

C4.1. Intermetallics. In MMCs, the presence of the reinforcement phases may affect the normal precipitation characteristics of intermetallic phases in comparison to those of the monolithic matrix alloy. The intermetallics also have their own characteristic electrochemical properties (e.g., corrosion potentials, pitting potentials, and corrosion current densities [58]). Noble intermetallics may induce galvanic corrosion of the matrix, whereas active intermetallics may go into dissolution and leave fissures or crevices. Hence, the difference between the corrosion characteristics of the MMC and its monolithic matrix alloy is also dependent on the extent to which the reinforcement phases affect the amount, distribution, and morphology of the intermetallics in the MMC. In Al/Al2O3 MMCs, Al8Mg5 and Mg2Si intermetallics provided corrosion paths along fiber–matrix interfaces [59]. Pits in Al/Al2O3 MMCs exposed to NaCl solutions containing H2O2 were attributed to the dissolution of MgAl3, which is rapidly attacked at low potentials [60]. In Al/mica MMCs, a dendritic phase which was probably Mg2Al3 or Al8Mg5 and spheroidized CuMgAl2 were preferentially attacked in nondeaerated 3.5 wt % NaCl [61]. C4.2. Dislocation Density. The high strength of particulate MMCs in comparison to their monolithic alloys is generated by high dislocation densities caused by a mismatch in the CTE between the reinforcement and matrix and heating and cooling histories [57]. Since cold working, which is the result of generating high dislocation densities, is known to change the corrosion behavior of metals such as steel [62] and aluminum [63], the corrosion behavior of MMCs may also be affected by high dislocation densities [31]. It has been suggested that corrosion near the SiC–Al interface in Al/SiC MMCs could be caused by high dislocation density due to a mismatch of the CTE between SiC and Al [64, 65]. C4.3. Physical Damage. MMCs can also be more susceptible to corrosion initiation at sites of physical damage in comparison to their monolithic alloys. Impacting the MMC

CORROSION OF MMC SYSTEMS

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matrix [66, 67], (2) microstructural chloride contamination originating from TiCl4 and BCl3 processing gases [68] that induced pitting [69] in some types of Al 6061/Gr/50 F-T6 MMCs in chloride-free environments, and extraneous carbon particles in Al 6092/B4C MMCs [70], exacerbating galvanic effects. Unlike the case with many types of metal alloys, MMCs cannot be always reprocessed to regain optimal properties. If an MMC had been subjected to high temperatures that lead to interphase formation, reprocessing the MMC to reverse the interphase formation would be unlikely. Also, fractured or damaged reinforcement constituents cannot be repaired. Therefore, it is very important to know the source and history of MMC products.

D. CORROSION OF MMC SYSTEMS The corrosion behavior of aluminum, magnesium, titanium, copper, stainless steel, lead, depleted uranium, and zinc MMCs is discussed. D1. Aluminum MMCs

FIGURE 35.8. Al–2 wt % Cu/Al2O3/F 50 MMC with localized deformation induced using a 1.58-mm-diameter Si3N4 ball (a) before and (b) after immersion in aerated 3.15 wt % NaCl at 30 C for five days. Notice (a) broken fibers and (b) enhanced corrosion initiation over deformed region. (Photos courtesy of Shruti Tiwari.)

surface may debond the matrix from the reinforcement or fracture the reinforcement (Fig. 35.8) with both cases creating microcrevices that can lead to corrosion initiation sites. In Al-2 wt % Cu/Al2O3/50F MMCs, regions of localized deformation corroded preferentially when the MMCs were exposed to aerated 3.15 wt % NaCl. Significant amounts of Cu precipitated on the surface of the deformed region, where corrosion was exacerabated. For similar levels of localized deformation, pure Al/Al2O3/50F MMCs were much less susceptible to preferred corrosion initiation at locally deformed regions. C4.4. Processing Deficiencies. Processing-induced corrosion is not inherently caused by the primary components of the MMC system but result from processing deficiencies such as too low or too high processing temperatures, contamination, and so on. Some cases that were observed in various MMCs were (1) low-integrity diffusion bonds that induced crevice corrosion between B MF and the aluminum

Aluminum is a reactive metal with a high driving force to revert back to its oxide, but it generally has good resistance to aqueous corrosion in near-neutral solutions due to the formation of a passive film [71]. In acidic and basic solutions, the passive film is not thermodynamically stable and thus corrosion rates are high [71]. Aluminum pits in halidecontaining solutions and the pitting potential (Epit) is linearly dependent on the logarithm of the halogen anion concentration [72]. However, in order for aluminum to pit in the opencircuit condition, it must be polarized to potentials noble to Epit by a cathodic reaction. Proton and oxygen reduction are two possible cathodic reactions, but in neutral, chloridecontaining solutions, oxygen reduction is necessary to initiate pitting. Pits do not nucleate on aluminum in the opencircuit condition if solutions are deaerated [73]. In aerated solutions, ultrapure aluminum (99.999 wt %), which is a poor catalyst for oxygen reduction, does not pit in the open-circuit condition. The slow oxygen reduction kinetics on ultrapure aluminum is believed to be caused by the high resistivity of aluminum oxide which restricts electron migration through the passive film [74, 75]. In aluminum alloys, however, noble precipitates and conducting reinforcements can polarize the alloy to Epit in aerated solutions. Corrosion of aluminum MMCs reinforced with boron, graphite, SiC, Al2O3, and mica will be discussed. Studies on stress corrosion and corrosion fatigue are also discussed. D1.1. Aluminum/Boron MMCs. Al/B/MF MMCs are usually fabricated by diffusion bonding B MFs between aluminum foils [76]. Aluminum borides (i.e., AlB12 and AlB2)

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have been found at B MF matrix interfaces [77]. Both the B MF ends and circumferential surfaces can support cathodic currents (Figs. 35.4 and 35.5) and therefore can induce galvanic corrosion. Hence, shielding the MFs from environmental exposure is a likely provision for corrosion control of these MMCs. D1.2. Aluminum/Graphite MMCs. Al/Gr MMCs are usually reinforced with either continuous or chopped fibers. The MMCs reinforced with continuous fibers have anisotropic properties with high specific tensile strength and stiffness along the fiber axis, but limitations of shear, compression, and transverse strengths generally excludes their use in structural applications [2]. The high thermal conductivity, negative coefficient of thermal expansion, and high stiffness of graphite, however, has made ultralow expansion Al/Gr MMCs ideal for thermally stable space structures. The advantages of graphite for enhancements in mechanical and physical properties, however, are tempered with negative characteristics regarding corrosion resistance that graphite imparts due to its low electrical resistivity, catalytic properties for oxygen reduction, and reactivity with aluminum at high temperatures. Severe galvanic corrosion is induced in Al/Gr MMCs in aerated solutions (Fig. 35.4), with the main cathodic reaction being oxygen rather than proton reduction [78]. At elevated temperatures, graphite reacts with aluminum, forming the Al4C3 interphase, which readily decomposes in water [46] to produce CH4 and aluminum hydroxide Al(OH)3 [51]. Also see Section C3. Al/Gr MMCs should therefore be carefully processed to avoid or minimize Al4C3 formation, and microstructures free of chlorides [69] should be ensured (see Section C4.4), either by careful control of processing parameters if chlorinated gases are used or using fabrication methods that do not use chlorinated gases [68]. Due to the reactivity of Al/Gr MMCs, they are more suited for use in dry environments. Certain corrosion-resistant procedures used for monolithic Al such as anodization may not be suitable for Al/Gr MMCs since the graphite fibers can be easily oxidized at positive potentials [79]. D1.3. Aluminum/Silicon Carbide MMCs. The Al/SiC MMCs have been reinforced with particles, whiskers, fibers, or monofilament. Due to the vast difference in electrical resistivities of SiC, variations in corrosion behavior may result. Hence, galvanic corrosion cannot be ruled out, and SiC can serve as an electrode for proton and oxygen reduction. The degree of galvanic corrosion is strongly dependent on the type of SiC reinforcement. The anodic polarization diagrams (Fig. 35.4) of ultrapure Al (99.999%), pure Al (99.9%), Al–2 wt % Cu, and Al 6061-T6 with the cathodic polarization diagrams of HP SiC, SiC MF (with either carbon cores or carbon-rich surface exposed) in aerated 3.15 wt %

NaCl shows that the galvanic current density of the aluminums coupled to SiC MF is approximately 15 times that when coupled to HP SiC of equal surface area. The influence of the carbon core and carbon-rich surface of the SiC MF is clearly seen (Fig. 35.4) where the polarization diagram of the SiC MF has a stronger resemblance to that of pitch-based graphite (P100 Gr) than that of HP SiC [25]. Galvanic current between a type of Nicalon SiC fiber and an aluminum alloy was also measured in an aerated NaCl solution [80], but the galvanic current was only 15% of that between carbon fiber and the aluminum alloy. Experimental results have generally indicated that the corrosion rate of particulate and whisker Al/SiC MMCs are higher than the monolithic matrix alloy [81–84] and increase with SiC content [84, 85] in aerated, chloride-containing environments. Weight loss data of 6092/SiC/P-T6 Al MMCs showed an increase in the corrosion rate as the SiC content increased from 5, 10, 20, 40, to 50 vol % for various 90-day humidity chamber tests [86]. At the 50 vol % SiC content level, the corrosion rate for an MMC with high-purity, highresistivity green SiC was noticeably lower than that of the MMC with low-resistivity black SiC [86]. The black SiC is likely to support more cathodic currents leading to higher corrosion rates in comparison to green SiC. The large variation in resistivity of the SiC may be the cause for conflicting results in the literature on the corrosion of Al/SiC MMCs. For example, no obvious evidence of galvanic corrosion was found in Al 6061/SiC MMCs with 17–27 vol % SiC particles [87]. The presence of SiC particles does not have a significant effect on the aluminum matrix passive current densities [84, 86, 88, 89] and pitting potentials [80, 86–88, 90–95]. Pit morphology, however, is indirectly affected by the presence of SiC particles. Pits on Al/SiC/W MMCs were notably more numerous and much smaller in size [96] compared to pits on wrought and powder-compacted monolithic alloys during anodic polarization in 0.1 N NaCl. The pits nucleated at intermetallic particles (not SiC), which are smaller and more numerous in the MMCs than in the monolithic matrix alloys [96]. SiC whiskers [96] and particles [87] can enhance the precipitation of the intermetallic phases. Pitting has also been observed at dendrite cores [97], near-eutectic silicon [94, 98], and intermetallic particles [98] in various Al/SiC MMCs. In the open-circuit condition, pits have also been observed to initiate at SiC–Al interfaces and could be caused by galvanic action with the SiC particle if the resistivity of SiC is relatively low or the hydrolysis of Al4C3 at the SiC–Al interface [43]. Since the corrosion behavior of Al/SiC MMCs is dependent on the microstructure, the corrosion characteristics can be altered by processing. In the MMC, void content [99], dislocation density [100], agglomeration of SiC particles [101], and the precipitation of active phases [102] are affected by processing conditions. Certain solution heat


treatments and high extrusion ratios improved the corrosion resistance of an Al 7091/SiC/20 P MMC [99]. Extrusion improved the corrosion resistance of cast MMCs by reducing the amount of pores and agglomerates of SiC particles [101]. Corrosion resistance was also improved by a finer, more homogenous distribution of secondary phases at the T4 temper in comparison to the O and F tempers [103]. It has been suggested that corrosion near the SiC–Al interface could be caused by high dislocation density due to a mismatch of the coefficient of thermal expansion between SiC and aluminum [25, 64, 65], segregation of alloying elements to the SiC–Al interface [104], or the formation of Al4C3, which hydrolyzes in water. Aluminum carbide has been identified as a source of corrosion for MMCs reinforced with particles [105] and SiC Nicalon fibers [80]. Also see Section C3. The formation of micro crevices caused by reinforcement particles left in relief as the matrix corrodes also exacerbates corrosion by localized acidification in anodic regions and alkalinization in cathodic regions in Al6092/SiC/P-T6 MMCs [45, 106]. The aluminum matrix loses its ability to passivate when the solution becomes either acidic or alkaline. Also See section C1.8. D1.4. Aluminum/Alumina MMCs. Particles and both short and continuous Al2O3 fibers have been used to reinforce aluminum alloys. Characteristic properties of Al/Al2O3 MMCs are low-weight, high-temperature tensile and fatigue strengths, low thermal conductivity and expansion, and superior wear resistance. Galvanic corrosion between Al2O3 and aluminum is not possible since Al2O3 is an insulator. The Al2O3 reinforcements usually do not have significant effects on pitting potentials [43, 80, 90, 92–94, 107] in chloride solutions. Passive current densities below the pitting potential have been reported to increase with Al2O3 content [107], although this may be related to processing since passive current densities for a particulate Al6092/ Al2O3/20P-T6 MMC were consistent with other MMCs in sodium sulfate solutions and in chloride solutions under the pitting potential [86]. Microbial corrosion was also reported to be more significant on particulate MMCs in comparison to the monolithic alloy, indicating that the Al2O3–Al matrix interface or Al2O3 particles may have aided biofilm formation [108]. In Al/Al2O3 MMCs, corrosion initiation usually occurs at intermetallic particles or contaminants introduced with the Al2O3 reinforcements. Hence, Al/Al2O3 MMCs with pure Al matrices usually have excellent corrosion resistance due to minimal amounts of intermetallic precipitates that can serve as cathodic sites [34]. Only slight corrosion damage was observed on pure Al/Al2O3/50 CF MMCs exposed to marine atmosphere 0.5 mile from the coastline for an 11-month period [79]. The presence of intermetallics and segregation of alloying elements may contribute to localized corrosion near

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reinforcements. Preferential corrosion near fibers [34, 60, 109, 110] and particles [94, 110, 111] is sometimes noticed in chloride-containing solutions. In a 2 wt % Mg aluminum alloy MMC, Fe and high levels of Mg (10 wt %) were detected near fibers [60]. It was suspected that the presence of Mg originated from Mg2Al3. Pitting near fibers was attributed to corrosion of Mg2Al3, which is rapidly attacked at low potentials [60]. The Al8Mg5 and Mg2Si intermetallics have also been reported to induce corrosion in Al/Al2O3 MMCs [59]. In Al–2 wt %Cu/Al2O3/50 CF MMCs, corrosion initiation occurred at copper-rich precipitates on the fiber– matrix interface [34]. In Al 6092/Al2O3/20 P-T6 MMCs, cathodic sites were identified as Fe–Si–Al intermetallics and low-resistivity Ti oxide or suboxide particles (likely introduced with the alumina reinforcement) [40]. D1.5. Aluminum/Mica MMCs. Muscovite (KAl3Si3O10 (OH)2) mica particles less than about 70 mm in size have been used in Al/mica MMCs [112] for potential use in applications where good antifriction, seizure resistance, and high-damping capacity are required [113]. Since muscovite is an insulator with resistivities that range from approximately 1013 to 1017 Wcm [27], galvanic corrosion should not be a problem. Muscovite is insoluble in cold water [114], but it has also been reported to absorb moisture and then swell [115]. Mica particles were cast in various aluminum alloys [115, 116]. In 3.5 wt % NaCl solution, the Al/mica MMCs had pitting potentials approximately 20–30 mV lower than the monolithic matrix alloys. In addition, intermetallics were preferentially attacked, regions around and away from mica particles pitted, mica–aluminum interfaces corroded, and mica particles exfoliated [50]. Also see Section C3. D1.6. Stress Corrosion Cracking of Al MMCs. There have been only a few studies on stress corrosion cracking and corrosion fatigue of DR and CR Al MMCs. D1.6.1 Discontinuous Reinforced MMCs. Stress corrosion cracking studies for alternate exposure and immersion in NaCl solutions have been conducted on aluminum MMCs reinforced with Al2O3 particles [117] and SiC particles [117–119] and whiskers [117]. The Al2024/Al2O3/P MMC was susceptible to stress corrosion cracking while subjected to three-point beam bending and alternate exposure or continuous immersions in a NaCl solution [117]. Under the same conditions, however, the 6061 Al MMCs reinforced with SiC particles and SiC whiskers were not susceptible to stress corrosion cracking [117]. Similarly, Al2024/SiC/P MMCs were not prone to stress corrosion cracking under constant strain at 75% of ultimate tensile strength while exposed to an aerated NaCl solution [118]. Slow strain rate tension testing of Al 2024/SiC/P MMCs indicated that the MMC lost up to 10% of failure strength compared to exposure in air [119].

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D1.6.2 Continuous Reinforced MMCs. Stress corrosion cracking studies for the immersed state have been conducted on aluminum MMCs reinforced with unidirectional, graphite fibers [120], boron monofilaments [66], and Nextel 440 (Al2O3, SiO2, B2O3) fibers [121]. Al6061/Gr/F MMCs were stressed parallel to the fiber axis in natural seawater. Failure was stress dependent at high stress levels and occurred in less than 100 h. At lower stresses, failure was primarily caused by extensive corrosion and therefore was relatively independent of stress levels. Al2024/B/MF MMCs stressed parallel to the fiber axis at 80% fracture strength in an NaCl solution did not fail in 1000 h but failed after 500 h when H2O2 was added to the NaCl solution. Extensive intergranular matrix corrosion and broken filaments at random sites were observed. The monolithic matrix alloy failed within 10 h under similar conditions. For Al2024/B/MF MMCs stressed perpendicular to the fiber axis at 90% yield strength in NaCl and NaCl with H2O2 solutions, failure occurred by intergranular matrix corrosion and separation at diffusion-bonded fiber–matrix interfaces. Failure times decreased with increasing B MF content; therefore, the presence of the MF was deleterious when stresses were perpendicular to the fiber axis. For the Al6061/Nextel/F MMCs, specimens were exposed to a pH 2 NaCl solution in the stressed and unstressed states [121]. The composite strength was measured before and after exposure to assess damage. The prevailing mode of failure was attributed to extensive corrosion along the fiber–matrix interface and not stress corrosion cracking. D1.7. Corrosion Fatigue in Al MMCs. Corrosion fatigue studies have been conducted on Al MMCs reinforced with graphite fibers [120], SiC whiskers [122–125], and SiC particles [123, 126]. Processing conditions and type of reinforcement affect corrosion fatigue behavior. Unnotched Al6061/Gr/F MMCs were exposed to natural seawater and stressed parallel to the fiber axis. The MMCs were processed with either silica (SiO2)–coated or TiB2-coated graphite fibers. For a given stress amplitude, the MMC with TiB2coated fibers had the longest corrosion fatigue life, followed by the MMC with the SiO2-coated fibers and the monolithic matrix alloy. At low stress amplitudes corresponding to longer exposure times, the MMC with the SiO2-coated fibers suffered premature failure due to extensive corrosion. In Al/ SiC MMCs, fatigue crack rates of compact tension specimens are usually higher in NaCl solutions as compared to air [122] or argon [126]. Loading frequency affects corrosion fatigue crack rates [126], but no consistent trends were observed. Fatigue [122] and corrosion fatigue [126] crack rates are influenced by loading and extrusion or rolling direction. The nucleation of a crack was also observed at the bottom of a corrosion pit [125]. The shape of the reinforcement constituent may also have significant effects on stress corrosion and corrosion fatigue, based on modeling that considers crack-tip strain rate [124]. The model predicts that crack rates are

reduced by increasing the reinforcement length-to-diameter ratio, which implies that MMCs reinforced with whiskers are more resistant to stress corrosion and corrosion fatigue than those reinforced with particles. This is in agreement with results on Al6061/SiC/W MMCs that were found to have longer corrosion fatigue lives than Al6061/SiC/P MMCs in salt-ladened moist air [123]. D2. Magnesium MMCs Magnesium is the lightest (density of only 1.7 g/cm3) and most active structural metal in the electromotive series [127]. Therefore, it has a very high driving force for corrosion, making it particularly susceptible to galvanic corrosion if it is coupled to noble reinforcement constituents. The normal corrosion of Mg is generally not affected significantly by dissolved oxygen [127] since the primary cathodic reaction in Mg corrosion is proton reduction. Hence, noble impurity elements that have low hydrogen overvoltage (e.g., iron, nickel, cobalt, and copper) [128] can significantly accelerate the corrosion rate of Mg, which is highly dependent on metallic purity [127]; for example, ultrapure Mg corrodes at the rate of 0.25 mm/y in seawater, but commercial Mg corrodes at about 100–500 times faster due to impurities [127]. If Mg is reinforced with constituents that are catalytic to oxygen reduction, MMC corrosion rates may significantly increase with aeration whereas that of the monolithic matrix alloy may not. Corrosion studies have been conducted on magnesium MMCs reinforced with B MF, graphite fibers, SiC MF, SiC particles, and Al2O3 fibers. The stress corrosion cracking behavior of an Mg/Al2O3 MMC has also been investigated. D2.1. Magnesium/Boron MMCs. Galvanic corrosion between Mg and pure boron is not a concern since pure boron is an insulator [129, 130]. However, tungsten core B MF is not an insulator due to the formation of tungsten boride [130]. Galvanic currents between virgin B MF (tungsten cores either shielded or exposed) and Mg [130] or an Mg alloy [129] were measurable in NaCl solutions. Galvanic currents were higher when tungsten cores were exposed [129, 130] since pure tungsten is an effective cathode [129]. Galvanic current densities increased approximately five times when Mg was coupled to B MF extracted from the matrix [129]. Corrosion rates of actual Mg alloy (MA2-1)/B/MF MMCs in 0.005 N and 0.5 N NaCl solutions were 12.5 and 81.7 g/m2 day, respectively, which were about six times the values of the monolithic matrix alloy in respective environments. D2.2. Magnesium/Graphite MMCs. The cathodic polarization diagrams of pitch-based graphite (cross section exposed) in aerated (Fig. 35.4) and deaerated (Fig. 35.5) 3.15 wt % NaCl with the anodic polarization diagrams of pure magnesium and Mg ZE41A in deaerated 3.15 wt %


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on the MMCs could have been caused by iron contamination of the magnesium matrix during processing in a steel crucible. Studies on a model MMC consisting of high-purity magnesium and well-separated SiC particles exposed to 3.5 wt % NaCl also did not show evidence of galvanic corrosion between the particles and matrix [135].

FIGURE 35.9. Exfoliated Mg AZ91C/Gr/12.7 F MMC with AZ31B Mg skins after a 20-year exposure period in an air-conditioned environment.

NaCl show that galvanic corrosion rates with graphite fiber will increase in aerated solutions. In addition, galvanic corrosion of magnesium is cathodically controlled, and therefore galvanic corrosion rates should increase with increasing area fraction of cathodic reinforcement material (Fig. 35.5). Actual MMCs immersed in air-exposed 0.001 N NaCl suffered severe degradation within five days [131]. Comparisons between a Mg AZ91C/Gr/40 P MMC and monolithic Mg AZ91C showed that the MMC open-circuit potential was approximately 0.3 V more noble and corrosion rate 40 times greater than that of the monolithic alloy in a deaerated 50-ppm chloride solution [132]. Even in relatively dry, air-conditioned environments an Mg AZ91C/Gr/12.7 F MMC with AZ31B Mg skins exfoliated over a 20-year period (Fig. 35.9). D2.3. Magnesium/Silicon Carbide MMCs. Galvanic corrosion between magnesium and SiC depends on the type of SiC reinforcement and on the presence of dissolved oxygen in solution. Anodic polarization diagrams of pure magnesium and Mg ZE41A with cathodic polarization diagrams of HP SiC and SiC MF exposed in aerated (Fig. 35.4) and deaerated 3.15 wt % (Fig. 35.5) show that galvanic corrosion rates (as determined by the mixed-potential theory) are greater in aerated solutions due to oxygen reduction [47, 133]. Galvanic corrosion rates are lower for couples with HP SiC as opposed to the SiC MF which have carbon cores and surfaces. Studies conducted on particulate Mg ZE41A alloy reinforced with 12 vol % SiC particles ranging in size up to approximately 20 mm did not show preferential attack between SiC particles and the matrix in salt spray tests [134]. Instead, macroscopic anodic and cathodic regions developed. Corrosion spread over the MMC surface much more rapidly than on the monolithic alloy, but the local corrosion rates were approximately only three times greater on the MMC. The authors [134] speculated that the higher corrosion rates

D2.4. Magnesium/Alumina MMCs. Galvanic corrosion should not be expected between magnesium and Al2O3 since Al2O3 is an insulator. The corrosion rates of an Mg AZ91C/ Al2O3/CF MMC [136] was approximately 100 times greater than that of the matrix alloy in 3.5 wt % NaCl at 25 C but similar to that of the matrix alloy in distilled water at 20 C. The significantly higher corrosion rates of the composites in the chloride solution appear to be caused by the presence of the Al2O3 fibers. Although galvanic corrosion is not expected between magnesium and Al2O3, conducting interphases or precipitates could potentially form due to the presence of the Al2O3 fibers. The open-circuit potential of a Mg AZ91C/ Al2O3/CF MMC was more noble than that of the matrix alloy in a 50-ppm Cl solution [137], indicating that noble precipitates or interphases could have been present. Stress corrosion tests [138] of Mg ZE41A/Al2O3/CF MMC in an NaCl–potassium chromate (K2CrO4) solution showed that notched and unnotched specimens stressed parallel to the fiber axis and exposed for approximately 100–1000 h in the NaCl–K2CrO4 solution retained approximately 90% of the strength in air. The matrix alloy and the MMC with the stress direction aligned perpendicular to the fiber axis retained only approximately 40–60% of the strength in air. D3. Titanium MMCs Titanium MMCs are being developed for aerospace, commercial, and biomedical applications. Titanium has a density of 4.5 g/cm3. D3.1. Titanium/Graphite MMCs. Porous titanium/Gr MMCs were processed and heat treated to fabricate a porous titanium/titanium carbide (TiC)/Gr MMC [17]. Polarization tests were conducted in 0.9 wt % NaCl and lactated Ringer’s solution for in vitro use. The anodic polarization current densities of the MMCs were significantly higher than that of pure monolithic titanium, which passivated. The authors attributed the higher corrosion rates of the Ti/TiC/Gr MMC to its porosity, which may have prevented complete passivation of the titanium matrix. Another possibility for the higher current densities could be the oxidation of graphite particles (see Section D3.2). D3.2. Titanium/Silicon Carbide MMCs. Corrosion studies on titanium alloy Ti–15V–3Cr–3Sn–3Al (Ti 15–3) [36] and titanium aluminide a2-Ti3Al (14 wt % aluminum, 21 wt %

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niobium, balance titanium) [139] reinforced with SiC MF have been conducted. The corrosion behavior of Ti 15-3/SiC/ MF MMC was investigated in 3.15 wt % NaCl. There was excellent agreement in the polarization diagrams of the actual MMC and that of a model utilizing the polarization diagrams of the individual constituents and the mixedpotential theory [36]. The matrix passivated and the carbon cores and carbon-rich outer surface of the SiC MF oxidized, likely forming CO2 similar to that observed in graphite fibers under anodic polarization [69]. In the open-circuit condition, the galvanic current density between Ti 15-3 and SiC MF cannot exceed that of the passive current density of Ti 15-3 based on their polarization diagrams (Fig. 35.4), and zeroresistance ammeter results confirmed that galvanic currents between Ti 15-3 and SiC MF were negligible. The corrosion behavior of the a2-Ti3Al/SiC/MF MMC [139] was somewhat similar to that of the Ti 15-3/SiC/MF MMC [36], with the exception that the a2-Ti3Al matrix is less resistant to pitting. During anodic polarization, the a2-Ti3Al/SiC/MF MMC pitted at approximately 1 VSCE in 0.5 N NaCl, which was approximately 0.5 V less than that of the monolithic matrix alloy. Some matrix pitting and crevice corrosion around the SiC MFs were also observed after anodic polarization. The galvanic current density of the a2-Ti3Al/SiC/MF MMC was negligible and limited to the passive current density of the a2-Ti3Al matrix [139]. D3.3. Titanium/Titanium Carbide and Titanium/Titanium Diboride MMCs. Particulate pure Ti/titanium carbide (pure Ti/TiC/P) and pure Ti/titanium diboride (pure Ti/TiB2/P) MMCs were fabricated by cold isostatic pressing following by sintering [140]. No interphase products were identified in the pure Ti/TiC/P MMC after processing, whereas TiB intephase products were identified in the Ti/TiB2/P MMC after processing. Composites reinforced with 2.5, 5, 10, and 20 vol % TiC or TiB2 were anodically polarized in deaerated 2 wt % HCl in the temperature range of 50–90 C. The passive current density for pure titanium was approximately 10 5 A/ cm2 throughout the temperature range. Generally, dissolution currents of the pure Ti/TiC/P and pure Ti/TiB2/P MMCs increased with increasing temperature and reinforcement content, and maximum values were about 20 times and 100 times, respectively, higher than that of pure titanium. For both types of MMCs, microscopy revealed that the titanium matrix was virtually uncorroded, whereas, the TiC particles underwent some degradation, and the TiB2 particles and TiB interphase were corroded significantly. D4. Copper MMCs Copper MMCs have been investigated for use in marine, electronic, and thermal applications. Copper is relatively heavy, with a density of 8.96 g/cm3. Reinforcements are typically chosen to impart strength and stiffness, reduce

weight, enhance thermal and electrical properties, improve machinability, and enhance wear resistance. Initial studies [1] were conducted on a wide variety of experimental copper and copper alloy MMCs reinforced with graphite, SiC, TiC, silicon nitride, boron carbide, and Al2O3 for marine applications. The MMCs generally showed corrosion behavior that was similar to that of the monolithic alloys, although corrosion rates were higher for some of the MMCs. Other studies have focused on copper MMCs for electronic, thermal, and tribological applications. D4.1. Copper/Graphite MMCs. The corrosion behavior of pure copper MMCs reinforced with 1.2–40 vol % graphite particles and 50 vol % graphite fibers in deaerated and aerated 3.5 wt % NaCl solutions [141] were investigated. The corrosion potential of the particulate-reinforced MMCs became more noble with increasing graphite content in both deaerated and aerated solutions, as would be expected by increasing the content of the noble graphite particles. The corrosion potential of the Cu/Gr/50 F MMC was approximately as noble as the Cu/Gr/40 P MMC in the aerated solution but was significantly more active than the particulate composite in the deaerated solution. This finding could be expected since oxygen reduction kinetics is normally diffusion limited in aerated solutions, but hydrogen evolution kinetics depend on the substrate and could be different on the graphite fibers as compared to on the graphite particles in deaerated solutions. The corrosion potentials of C90300 copper alloy (with 1 wt % titanium additive to increase graphite wettability) MMCs reinforced with 1–10 vol % graphite particles also increased with increasing graphite content in an aqueous solution containing ferric chloride, copper sulfate, and hydrochloric acid [12]. D4.2. Copper/Silicon Carbide MMCs. The corrosion behavior of pure copper MMCs reinforced with 0, 5, 10, and 20 vol % SiC particles were examined in a 5 wt % NaCl solution [142]. Porosity in the materials ranged from 2.2 to 3.5% and generally increased with increasing SiC content. Corrosion potentials became more active, and corrosion current densities increased with increasing SiC content. Decreasing corrosion potentials would not be expected with increasing SiC content if SiC served as an efficient cathode. The corrosion morphology indicated that there was significant corrosion at SiC–copper interfaces. Voids caused by porosity and SiC–copper interfaces both increased with increasing SiC content. Hence, the decrease in corrosion potential with increasing SiC content is likely to have been caused by an increase in anodic sites at voids and SiC–copper interfaces. D4.3. Copper/Alumina MMCs. The corrosion behavior of copper MMCs reinforced with 2.7 vol % Al2O3 was


examined in deaerated and aerated 3.5 wt % NaCl [143]. Galvanic corrosion with Al2O3 is not expected since Al2O3 is an insulator. The corrosion rates of the MMCs were comparable to that of monolithic copper, and the corrosion potentials of the MMC were only 0.01–0.02 V, more active than that of monolithic pure copper. D5. Stainless Steel MMCs Sintered, particulate composites consisting of ferritic 434L stainless steel (SS) and Al2O3 particles have been developed for potential application in chemical processing plants, turbine blades, and heat exchanger tubes [144–146]. Austenitic 316L SS reinforced with Al2O3 and Y2O3 have also been investigated for enhanced strength and wear resistance [147]. D5.1. Stainless Steel/Alumina MMCs. The corrosion behavior of sintered 434L SS/Al2O3 MMCs and sintered 434L SS alloy without Al2O3 particles was examined [144–146]. The volume percent of Al2O3 particles in these materials ranged from 0 to 8%. The effect of small amounts of titanium and niobium alloying elements on corrosion resistance was also investigated. Galvanic corrosion between 434L SS and Al2O3 should not occur since the latter is an insulator. In 1 N H2SO4 [144–146], there was no strong correlation between Al2O3 content and corrosion behavior. One of the few generalities that could be made was that passive-current densities were high and within an order of magnitude of 1 mA/cm2 for almost all materials. In the 5 wt % NaCl solutions, iCORR of the MMCs was less than 10 mA/cm2 [145]. Upon polarization, all materials displayed active corrosion behavior in the NaCl solutions. Particulate 316L SS MMCs [147], fabricated using powder metallurgy, were reinforced with 3, 4, and 5 wt % Al2O3 and additions of 2 wt % chromium diboride (CrB2) or 1 wt % boron nitride (BN) for sintering aids. The density of the MMCs ranged from 86 to 96% of the theoretical value. Unreinforced 316L SS specimens were also fabricated using powder metallurgy without sintering aids, resulting in 85% of theoretical density. Less porosity was present in the reinforced MMCs as compared to the unreinforced pure 316L SS specimen. The test samples were immersed in 10 wt % sulfuric acid (H2SO4) at room temperature for 24 h, 1 wt % hydrochloric acid (HCl) at room temperature for 24 h, and boiling 10 wt % nitric acid (HNO3) for 8 h. The pure, unreinforced 316L SS specimens passivated in the 10 wt % H2SO4 solution, whereas the corrosion rate of the MMC generally increased with increasing Al2O3 content to a maximum value of approximately 4 mm/yr. The MMCs performed better than the unreinforced 316L SS specimen in the 1 wt % HCl solution but worse than the unreinforced specimen in the boiling nitric acid solution. There was no strong correlation between Al2O3 content in the MMCs and

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the corrosion rates in 1 wt % HCl and boiling 10 wt % HNO3 solutions. D5.2. Stainless Steel/Yttria MMCs. Yttria (Y2O3) is an insulator and galvanic effects are not expected. The 316L SS specimens discussed above [147] were also reinforced with 3, 4, and 5 wt % Y2O3 and additions of 2 wt % chromium diboride (CrB2) or 1 wt % BN as sintering aids. In all solutions (i.e., sulfuric, hydrochloric, and nitric acid solutions), the Y2O3-reinforced MMCs exhibited reduced corrosion resistance as compare to the Al2O3-reinforced MMCs. The Y2O3 MMCs were sintered to 88–96% of theoretical density, and the Al2O3 MMCs were sintered to 86–92% of theoretical density. The Y2O3 particles also showed better bonding to the matrix, probably forming a complex YCrO3 oxide, as compared to the Al2O3 particles. It is possible that the formation of the reaction layer around the Y2O3 particles may have depleted chromium from the matrix, resulting in reduced corrosion resistance, as compared to the Al2O3reinforced MMCs. D6. Lead MMCs Lead is a relatively heavy metal with a density of 11.4 g/cm3. Lead MMCs, therefore, are normally developed for applications where a combination of its structural, physical, and chemical properties is important. The corrosion behavior of pure lead MMCs in simulated lead–acid battery environments has been studied to assess the feasibility of using these composites as positive electrode grids in place of conventional lead-based alloy grid materials. Lead can be alloyed with elements such as arsenic, antimony, or calcium to increase strength and stiffness. These elements, however, reduce corrosion resistance. Monolithic pure lead has very good corrosion resistance in lead–acid battery environments (which consists of sulfuric acid solutions) but is heavy and lacks sufficient mechanical strength. Pure lead, therefore, has been reinforced with strong, lightweight fibers in hopes of achieving the goals of increasing strength, reducing weight, and retaining the corrosion resistance of pure lead [148–150]. For other applications, discontinuous reinforced lead– antimony alloy MMCs were also studied in sodium chloride solutions [151]. To simulate corrosion in lead–acid battery environments, lead MMC reinforced with Al2O3, carbon, SiC, and glass– quartz fibers of various volume percents have been [148– 150] anodically polarized at 1.226 V (vs. mercury/mercurous sulfate reference electrode) in sulfuric acid solutions (of 1.285 specific gravity) at 50, 60, and/or 70 C. At 1.226 V, lead and water are oxidized to lead dioxide (PbO2) and molecular oxygen (O2), respectively [152, 153]. About one-third of the total anodic current is consumed in the oxidation of lead under these conditions [150]. Poor bonding between Al2O3 fibers and the matrix allowed the electrolyte

496


to diffuse into fiber–matrix interfaces, leading to accelerated corrosion [148] and swelling of the composite due to corrosion product buildup [150]. The graphite fibers were also subjected to oxidation [149]. Lead (80 wt %)–antimony (20 wt %) alloy MMCs reinforced with 1–5 wt % zircon (ZrSiO4) particles [151] were exposed to a 1 N NaCl solution. Zircon should not induce galvanic corrosion. Weight loss measurements, made over a 72-h period, showed that the corrosion rate of the MMCs increased with increasing zircon content. D7. Depleted Uranium MMCs Depleted uranium/tungsten fiber (DU/W/F) MMCs are the antithesis of the lightweight MMCs and were developed to create high-density materials. Uranium has a density of 18.9 g/cm3. DU corrodes galvanically when coupled to tungsten fibers in air-exposed 3.5 wt % NaCl solutions at room temperature [154]. The open-circuit potential of tungsten fiber ( 0.25 VSCE) is noble to that of the DU alloy ( 0.80 VSCE). The open-circuit potentials of the DU/W/F MMC and galvanic couples consisting of tungsten fiber and DU alloy of equal areas are 0.78 and 0.77 VSCE, respectively, and fall between those of tungsten fibers and the DU alloy. The galvanic corrosion current density measured between equal areas of tungsten fibers and the DU alloy was equal to about 4 10 5 A/cm2. In a 30-day exposure test in the NaCl solution, the DU/W/F MMC lost 43.56 mg/cm2, which was about 1.3 times that of the DU alloy. D8. Zinc MMCs Zinc MMCs have been developed [155] for potential use as bearing materials. Zinc has a density of 7.14 g/cm3. Zinc alloy ZA-27 MMCs were cast with 1, 3, and 5 wt % graphite particles ranging in sizes from 100 to 150 mm. Zinc alloys are known to have excellent wear and bearing characteristics [156]. The zinc MMCs were resistant to corrosion in SAE 40 grade lubricant that had been in service for six months in an internal combustion engine. In 1 N HCl, the corrosion rates of the MMCs decreased with time.

E. CORROSION PROTECTION OF MMCs Corrosion of metals can be prevented with the use of protective coatings and inhibitors. The use of impervious, inhibitive, or cathodically protective coatings will depend on the application and substrate. Selecting suitable coatings for MMCs will likely require testing and verification because a proven coating system for an alloy may not be effective for an MMC of that alloy. Poor adhesion and wettability between the coating and reinforcement or differences in the

electrochemical properties of the alloy and MMC may render a good coating system for the alloy ineffective for the MMC. Other coating techniques such as anodization could also be ineffective or even deleterious to the MMC. When an aluminum MMC is anodized, for example, the reinforcements can impede the growth of a continuous aluminum oxide film [157] or the reinforcements can be compromised by oxidation such as in the case of graphite fibers that oxidize to CO2 [46]. The use of inhibitors is usually reserved for closed systems and, therefore, may not be an option in many practical cases. An inhibitor intended for a monolithic alloy should not be used for an MMC of that alloy until ample examination confirms its effectiveness. Various studies on the corrosion protection of MMCs utilizing organic coatings, inorganic coatings, anodization, and chemical conversion coatings have been summarized elsewhere [158], and results are given in Table 35.2. The studies have generally shown that the best protection for MMCs that are susceptible to corrosion has been achieved by completely shielding the MMC from the environment utilizing coatings. Undoubtedly, in the future, generally acceptable methods to protect MMCs will likely be developed.

F. SUMMARY There are many additional concerns regarding the corrosion of MMCs in comparison to their monolithic matrix alloys. Certain MMC systems have inherent corrosion problems. Galvanic corrosion of the matrix may be induced by conductive reinforcements (e.g., Gr, some semiconductors with high impurity levels, and other metals). If reinforcements are semiconductive, n types could suppress galvanic corrosion of the matrix under illumination, while p types could induce galvanic corrosion under illumination. When the reinforcements are insulators, galvanic corrosion of the matrix will not be induced; however, the corrosion behavior of the MMC in comparison to the monolithic matrix alloy may be different. Whether the reinforcements are conductive, semiconductive, or insulating, they can affect the precipitation of intermetallics, induce higher dislocation densities, and react with the matrix-forming interphases. All of these phenomena may alter the corrosion behavior of the MMC in comparison to the monolithic matrix alloy. Corrosion initiation in MMCs may also be more sensitive to physical damage. Impacting the MMC surface may lead to debonding at the reinforcement– matrix interface or fracturing the reinforcement leading to the formation of minute crevices, which may enhance corrosion initiation. The problem could possibly be exacerbated with MMCs having higher strength matrices and higher reinforcement content. Caution should be used in assuming that MMCs belonging to a specific MMC group (e.g., Al/SiC MMCs) all have similar corrosion behavior since variations

NOMENCLATURE

TABLE 35.2.

Summary of Corrosion Protection Studies on MMCs

MMC Type

Substrate Protected

Al/Gr

Al/SiC

Al/Al2O3

Coating/Treatment

Environment

Outcome

References

Organic coatings

Marine and NaCl solution NaCl solution Marine

Protection

[159–161]

Short-term protection No protection

[162] [159]

Marine

Protection

[159]

Marine

Accelerated corrosion

[163, 164]

Marine

Delamination from exposed edges Protection if panel edges were sealed to prevent Al/Gr exposure Protection if panel edges sealed Protection if edges sealed with epoxy Delayed pitting on surface foils Protection

[159]

Protection Protection Various levels or protection Limited protection

[163] [163] [167–170]

Limited protection

[172, 173]

Surface Al foils on MMC Al/Gr MMC Surface Al foils on MMC Surface Al foils on MMC Surface Al foils on MMC Al/Gr MMC

Inorganic diamondlike coating CVD and PVD inorganic coatings Electroplated Ni coating (without defects) Electroless Ni coating (with defects) Ti and Ni cladding

Surface Al foils on MMC

Electrodeposited Al/Mn on electroless Ni coating

Marine

Surface Al foils on MMC Surface Al foils on MMC Surface Al foils on MMC Al/SiC MMC

Anodization with dichromate sealing Chromate/phosphate conversion coating Chemical passivation with CeCl3 Organic epoxy coating

Marine

Al/SiC MMC Al/SiC MMC Al/SiC MMC

Plasma-sprayed alumina coating Flame-sprayed Al coating Anodization

Al/SiC MMC

Chemical passivation with CeCl3 Chemical passivation with CeCl3

Al/Al2O3 MMC

in the quality of the reinforcement and matrix alloy, manufacturing technique (e.g., powder metallurgy versus casting), thermomechanical processing, and other factors can all alter corrosion behavior. Hence, it will be difficult to obtain generalized and consistent corrosion behavior for a specific MMC group until standards are developed for the manufacture and processing of MMCs.

Marine NaCl solutions Marine and NaCl solution Marine Marine Marine and NaCl solution NaCl solution NaCl solution

HP I ic iCORR iGALV MMC SC SiC MF SiCMFE SiCMFS

NOMENCLATURE Superscript E B MF BMFE BMFS CD Epit EGALV Gr GrE

497

Boron monofilament BMF electrode with MF ends exposed BMF electrode with MF circumferential surface exposed Current density Pitting potential Galvanic couple potential Graphite Gr electrode with fiber ends exposed

Superscript S T vol % VSCE wt % xC or XC r

[159]

[163] [163] [165, 166] [160, 163]

[165, 171]

Hot pressed Current density Cathodic current density Corrosion current density Galvanic current density Metal–matrix composite Semiconductor Silicon carbide monofilament SiCMF electrode with MF ends exposed SiCMF electrode with MF circumferential surface exposed Electrode with fiber or MF ends exposed (e.g., SiCMFE) Electrode with fiber or MF circumferential surface exposed (e.g., SiCMFS) Thickness Volume percent Volts versus a calomel electrode Weight percent Cathodic area fraction Resistivity

498


Units A cm h s V W

Amperes Centimeter Hours Seconds Volts Ohm

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97. G. E. Kiourtsidis and S. M. Skolianos, “Corrosion Behavior of Squeeze-Cast Silicon Carbide-2024 Composites in Aerated 3.5 wt.% Sodium Chloride,” Mater. Sci. Eng. A, 248, 165–172 (1998). 98. M. M. Buarzaiga and S. J. Thorpe, “Corrosion Behavior of AsCast, Silicon Carbide Particulate-Aluminum Alloy MetalMatrix Composites,” Corrosion, 50(3), 176–185 (1994). 99. R. C. Paciej and V. S. Agarwala, Corrosion, 44, 680–684 (1988). 100. S. R. Nutt and J. M. Duva, Scripta Metallurgica, 20, 1055–1058 (1986). 101. M. S. Bhat, M. K. Surappa, and H. V. Sudhaker Nayak, “Corrosion Behaviour of Silicon Carbide Particle Reinforced 6061/A1 Alloy Composites,” J. Mater. Sci., 26(18), 4991–4996 (1991). 102. J. England and I. W. Hall, Scripta Metallurgica, 20, 697–700 (1986). 103. Z. Ahmad and B. J. Abdul Aleem, “Effect of Temper on Seawater Corrosion of an Aluminum-Silicon Carbide Composite Alloy,” Corrosion, 52(11), 857–864 (1996). 104. W. N. C. Garrard, “The Corrosion Behaviour of AluminumSilicon Carbide Composites in Aerated 3.5% Sodium Chloride,” Corros. Sci., 36(5), 837–851 (1994). 105. J. K. Park and J. P. Lucas, “Moisture Effect on SiCp/6061 A1 MMC: Dissolution of Interfacial A14C3,” Scripta Mater., 37 (4), 511–516 (1997). 106. L. H. Hihara, H. Ding, and T. Devarajan, “Corrosion-Initiation Sites on Aluminum Metal-Matrix Composites,” U.S. Army Corrosion Summit 2004, Cocoa Beach, FL, 2004. 107. C.-K. Fang, C. C. Huang, and T. H. Chuang, “Synergistic Effects of Wear and Corrosion for A12O3 ParticulateReinforced 6061 Aluminum Matrix Composites,” Metall. Mater. Trans. A, 30A, 643–651 (1999). 108. R. U. Vaidya et al., “Effect of Microbiologically Influenced Corrosion on the Tensile Stress-Strain Response of Aluminum and Alumina-Particle Reinforced Aluminum Composite,” Corrosion, 53(2), 136–141 (1997). 109. V. S. Agarwala, Abstract No. 15 in Extended Abstracts, The Electrochemical Society, Montreal, Canada, May 1982. 110. L. Bertolini, M. F. Brunella, and S. Candiani, “Corrosion Behavior of a Particulate Metal-Matrix Composite,” Corrosion, 55(4), 422–431 (1999). 111. J. M. G. DeSalazar et al., “Corrosion Behaviour of AA6061 and AA7005 Reinforced with Al2O3 Particles in Aerated 3.5% Chloride Solutions: Potentiodynamic Measurements and Microstructure Evaluation,” Corros. Sci., 41, 529–545 (1999). 112. D. Nath, R. T. Bhat, and P. K. Rohatgi, J. Mater. Sci., 15, 1241–1251 (1980). 113. P. K. Rohatgi, R. Asthana, and S. Das, Int. Mater. Rev., 31, 115 (1986). 114. R. C. Weast, CRC Handbook of Chemistry and Physics, 67th ed., CRC Press, Boca Raton, FL, 1986, p. B-116. 115. D. Nath and T. K. Namboodhiri, Composites, 19, 237–243 (1988).

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36 ENVIRONMENTAL DEGRADATION OF ENGINEERED BARRIER MATERIALS IN NUCLEAR WASTE REPOSITORIES R. B. REBAK GE Global Research, Niskayuna, New York

A. B. C. D. E.

Introduction National programs Environments and materials Degradation mechanisms Reducing or anoxic environments E1. Carbon steel and low-alloy steel E2. Copper E3. Stainless steel and nickel alloys E4. Titanium F. Oxidizing environments F1. Corrosion behavior of alloy 22 (N06022) F2. Uniform and passive corrosion of alloy 22 F3. Localized corrosion of alloy 22 F4. Environmentally assisted cracking of alloy 22 F5. Corrosion behavior of titanium alloys G. Summary and conclusions References

A. INTRODUCTION All the countries addressing the issue of nuclear waste, mainly produced by nuclear power plants, are considering disposing them in stable geologic repositories. In most of the repositories (e.g., Finland, Sweden, Canada, France) the environments will be reducing in nature, except for the repository in the United States, in which the ingress of oxygen will not be restricted. For the reducing repositories the different national programs are considering materials as

carbon steel and copper. For the repository in the United States, some of the most corrosion-resistant commercially available alloys, such as nickel and titanium alloys, are being characterized. This chapter presents a summary of the behavior of the different materials under consideration for the repositories and the current understanding of the degradation modes of the proposed alloys in groundwater environments from the point of view of general corrosion, localized corrosion, and environmentally assisted cracking. Recent concerns about global warming and the release of greenhouse gases by the fossil fuel power industry have reignited the consideration of alternative sources of energy such as wind, solar, fuel cells, and nuclear power. Currently there are more than 40 nuclear power reactors under construction in the world, most of them in Asia [1]. Because of this, some reports claim that there is a nuclear power renaissance. However, even though nuclear power has been used for more than 60 years, the issue of the toxic radioactive waste generated during energy production still needs to be resolved. Radioactive materials are pertinent not just to nuclear power since they are also used worldwide in other fields, including medical applications and production of weapons. Once the radioactive materials lose their commercial value, they are considered radioactive waste, and they need to be isolated from the environment until the radioactive decay has reduced its toxicity to innocuous levels for plants, animals, and humans. Different types of radioactive waste are produced during commercial and defense nuclear fuel cycles. One type of waste, denoted high-level waste (HLW), contains the highest concentration of radiotoxic and heat-generating species. Because of this factor, the most stringent standards for disposing of radioactive wastes are being placed


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worldwide on HLW, and the majority of the radioactive waste management effort is being directed toward the HLW problem. One of the most common and most voluminous types of HLW is the spent fuel (SF) from commercial nuclear reactors for power generation. All of the countries currently studying the options for disposing of HLW have selected deep stable geologic formations to be the primary barrier for accomplishing this isolation. It is postulated that, by the very nature of these geologic sites, they will contain the waste for long times, limiting their spread, for example, through water flow. All the repository designs also plan to delay the release of radionuclides to the environment by the construction of engineered barrier systems (EBS) between the waste and the geologic formation. These barriers will be installed to limit water reaching the repository and to restrict radionuclide migration from the waste. The principal engineered component in this multibarrier approach is the waste package, which includes the waste itself, possibly a stabilizing matrix for the waste such a glass, and a metallic container that encloses the wasteform. Beyond the metallic containers, other secondary barriers could be added to attenuate the impact of the emplacement environment on the containers. The secondary barriers may include a drip shield such as in the U.S. design or backfilling with bentonite such as in the Canadian and other designs (Table 36.1) [2, 3]. A discussion in detail of the characteristics and significance of the engineered barrier systems in all the planned repositories is given by Bennett et al. [4] More than 30 nations are currently considering the geologic disposal of HLW [3–6]. A short list of these nations is in Table 36.1. Twenty years ago most of the repository designs specified lifetimes from 300 to 1000 years. Currently, the minimum length of time specified for some repositories has increased to 10,000, 100,000, and even 1,000,000 years [7–11]. The viability of extrapolating degradation data from short-term testing to long-time performance has been addressed by some investigators and the American Society for Testing and Materials (ASTM) [7, 12, 13]. Others have

proposed models to predict the lifetime performance of container alloys [14, 15].

B. NATIONAL PROGRAMS Table 36.1 lists some countries that are currently considering geologic repositories for nuclear waste. One of the most advanced studies for a repository corresponds to the United States, which is planning to locate its nuclear waste at a remote desert site in Nevada [16]. The container for the waste will be a double-walled cylinder having a 2.5-cm-thick layer of alloy N06022 on the outside and a 5-cm-thick layer of nuclear-grade type 316 stainless steel in the inside. After more than two decades of scientific investigations, the U.S. Department of Energy submitted a license application to build the repository on June 3, 2008, and the Nuclear Regulatory Commission accepted this application on September 8, 2008 [11, 17]. After a formal review process of three to four years, construction may start in late 2011, and the first waste emplacement may not occur until 2020 [7, 18]. Currently, the Yucca Mountain repository project does not seem a high priority of the U.S. Department of Energy [18]. The United Kingdom has also recognized that the geologic disposal of the waste in a mined repository is the best available approach [19]. In June 2008 the U.K. government issued the white paper “Managing Radioactive Waste Safely,” where a framework for implementing geologic disposal is outlined. The location for the repository will be defined through geologic screening and community engagement. To complement the permanent repository studies, the U.K. Committee for Radioactive Waste Management has also recommended a robust program of interim storage [19]. The Finnish repository will be located in crystalline bedrock at Olkiluoto Island on the western coast of Finland. The waste containers will be made using nodular cast iron with a 50-mm-thick overpack of copper [20]. The repository in Finland should start operations in 2020 and will continue for

TABLE 36.1. Selected Countries Considering Repositories for High-Level Waste Country

Possible Environment, Host Rock

Scheduled Start Operations

Container Materials Being Studied

Belgium Canada Finland France Germany Japan

Reducing or anoxic, clay Reducing or anoxic, granite Reducing or anoxic, granite Reducing or anoxic, clay Salt dome layer Reducing or anoxic, granite plus bentonite buffer Reducing or anoxic, granite plus clay Reducing or anoxic, clay Oxidizing, non saturated, volcanic tuff

2035 — 2020 — 2030 Late 2030s

Carbon steel, cement, stainless steel Carbon steel insert, copper Cast iron insert, copper Carbon steel Carbon steel Carbon steel, titanium

2023 — Initially 1998, now 2020

Cast iron insert, copper Carbon steel Ni–Cr–Mo Alloy C-22, titanium Gr 7, 28, and 29

Sweden Switzerland United States

ENVIRONMENTS AND MATERIALS

approximately 100 years [21]. Sweden has elected a site for its underground repository in the municipality of € Osthammar, 500 m below ground in crystalline rock [22, 23]. The waste will be packed in cast iron baskets inside thick copper canisters surrounded by bentonite clay. Each container is a double-walled cylinder of approximately 1 m diameter and 5 m long [22]. The Swedish repository is scheduled to open in 2023 and it is designed to contain the waste for 100,000 years [22]. The Japanese Final Disposal Plan calls for a repository that will start operating in the late 2030s [24]. The final site for the repository has not been selected yet, but two underground research laboratories have been selected, one 1000 m deep in crystalline rock in the presence of fresh water and the second in sedimentary rock 500 m deep in the presence of saline water [25]. In the final Japanese repository the metal containers will be surrounded by bentonite buffer material. The final material for the container has not been selected yet, but it is reported that a thick steel container surrounded by a bentonite buffer overpack would be robust design [24]. The waste disposal for the French nuclear industry has been outlined in the document Dossier 2005 and calls for the commissioning of a disposal facility by the year 2025 [26, 27]. An important concept in the design of the French repository is its reversibility or design evolution at all steps for at least 100 years [26]. The cylindrical containers for the high-level waste (type C in a glass matrix) will be made of standard steel 5 cm thick, 60 cm diameter, and approximately 1.5 m long. It is estimated that this container will remain leak proof for 4000 years [26]. The spent-fuel container may have a wall thickness of over 10 cm and last 10,000 years. Lithuania has recently announced that it will follow the French design for its own repository. Germany is exploring the possibility of a repository in the salt dome at Gorleben probably using steel containers for the nuclear waste. The waste repository should be stable for a million years and the containers should be retrievable for the entire operation time.

C. ENVIRONMENTS AND MATERIALS Table 36.1 gives a list of countries that are planning to build repositories and the general characteristics of each repository. As mentioned before, each repository consists of a stable geologic formation with the addition of engineered barriers (e.g., container). The repositories can be divided into two large groups according to the nature of the environment, oxidizing (Yucca Mountain, U.S.) and reducing (the other nations). For the reducing (no-oxygen) repositories, the most common host materials are clay, basalt, salt, and granite [2]. The containers are intended to be placed in alcoves located at varying depths below the water table. The depth of emplacement may vary from country to country, but it is generally assumed to be of the order of 500 m. The United States has the

505

only nonsaturated (above-the-water-table) repository design, with unrestricted access of oxygen [9]. According to the value of the redox potential, the environment of the world repositories can be categorized as reducing or oxidizing. Under reducing conditions, the cathodic reaction is controlled by the hydrogen evolution reaction. On the other hand, oxidizing conditions are characterized by cathodic reactions such as the reduction of dissolved oxygen. Most of the repositories in the world will be reducing based on redox potentials, since they will rely on depth (where the solubility of oxygen in water is minimal) and a projected backfill with bentonite [3]. One of the intended functions of the backfill is to retard the diffusion of any available oxygen toward the containers and the diffusion of the radionuclides away from the containers. The repository in the United States will not have restrictions regarding the availability of oxygen to contact the containers, that is, the redox potential will be oxidizing in nature, provided an aqueous solution materializes. The groundwaters associated with the rock formations should all be relatively benign to most materials because of their low ionic strengths, near-neutral pH, and low concentrations of halide ions [3, 6]. The corrosivity of the groundwaters could increase if significant vaporization occurs when the containers experience higher temperatures during the early emplacement times. The container temperature will be influenced by the design and loading of the waste package, the density of waste package emplacement, and the thermal properties of the surrounding rock. Because heat is a significant byproduct of HLW decay, the temperature of the waste containers will initially increase and then decrease as the activity of the waste decays. The predicted maximum temperature for waste packages emplaced in a consolidated volcanic ash (tuff) formation in the United States is not expected to be higher than 160–200 C [9]. Typical maximum container temperatures for a number of other repository locations are expected to be lower than 100 C [3, 4]. The International Atomic Energy Agency (IAEA) offers guidelines for the minimum requirements such as tests and evaluations that the waste packages (containers) should undergo during the selection process [28]. The degradation mechanisms that need to be studied include corrosion, microbial activity, and radiation damage. Table 36.1 shows that, except for the United States, most of the recommended materials for the containers will be carbon steel, stainless steel, or copper [2, 4]. Since for most of the nations the environment is rather benign or controlled, the alloys selected for the containers are not in the high end of the scale of the corrosion-resistant alloys. For the U.S. containers, some of the most corrosion-resistant materials currently available have been recommended, including the Ni–Cr–Mo alloy 22 and titanium grade 7. The compositions of some of the candidate materials for the engineered barriers being studied worldwide at this moment are given in Table 36.2.

506


TABLE 36.2. Approximate Chemical Composition (wt %) for Candidate Alloys Alloy

UNS

ASTM

Gray cast iron

F10001–F10012 A319–A159

1018 Carbon steel 4130 Alloy steel 2.25Cr–1Mo Type 304 Type 316 Copper Monel 400 Incoloy 825 Inconel 625 Hastelloy C-4 Hastelloy C-22 Ti Gr 2 Ti Gr 7 Ti Gr 16 Ti Gr 12 Ti Gr 29

G10180 G41300 K30736 S30400 S31600 Various N04400 N08825 N06625 N06455 N06022 R50400 R52400 R52402 R53400 R56404

A29 A29 A213 A182 A182 — B127 B163 B366 B575 B575 B265 B265 B265 B265 B265

Cr

Cu

Fe

Mo

Ni

Ti

Other

—

—

95 (bal)

—

—

—

98 (bal) 97 (bal) bal 70 (bal) 67 (bal)

— 0.2 1 — 2.5

— — — 9 12

— — — — —

3–3.5 C, 2–2.4 Si, 0.8 Mn 0.18 C, 0.5 Mn 0.3 C, 0.5 Mn 0.05 C, 0.4 Mn, 0.2 V 2 max Mn, 1 max Si 2 max Mn, 1 max Si

2.5 max 30 (bal) 5 max 3 max 4 0.3 max 0.3 max 0.3 max 0.3 max 0.25 max

— 3.0 9.0 16 13 — — — 0.3 —

66.5 42 60 (bal) 65 (bal) 57 (bal) — — — 0.8 —

— 0.9 0.2 — — 99 (bal) 98 (bal) 98 (bal) 98 (bal) 90 (bal)

2 max Mn 1 max Mn, 0.5 max Si 4 Nb, 0.5 max Mn 2 max Co 3 W, 2.5 max Co 0.25 max O 0.2 Pd, 0.25 max O 0.06 Pd, 0.25 max O 0.25 max O 6Al, 4V, 0.08-0.14 Ru

— — 1.0 — 2.25 — 19 — 17 — — 99.9 — 32 (bal) 21.5 2.2 21.5 — 16 — 22 — — — — — — — — — — —

D. DEGRADATION MECHANISMS The materials that are being characterized for their corrosion resistance behavior as engineered barriers applications are listed in Tables 36.1 and 36.2. It could be anticipated that most of these materials would suffer several types of corrosion processes, which in general can be grouped as (1) general corrosion, (2) localized corrosion, and (3) environmentally assisted cracking or stress corrosion cracking (SCC). One mode of corrosion that has been extensively investigated in several national programs is microbiologically influenced corrosion (MIC). MIC could affect the three corrosion processes mentioned above mainly by changing the environment in the vicinity of the containers. King provided a decision tree approach to determine if MIC would be an important factor in determining the lifetime of the containers [29]. Figure 36.1 shows schematically how metallurgical and environmental factors may control the occurrence of the three main corrosion processes. If water is present in the repository, the container material would adopt a characteristic potential called the rest potential or corrosion potential (Ecorr). Localized corrosion such as crevice corrosion or pitting corrosion may happen only above a threshold potential or critical potential (Ecrit). That is, if the material of the container adopts a potential that is below Ecrit, only general or uniform corrosion may occur. Environmentally assisted cracking (EAC) may occur at any potential; however, for EAC to occur, the simultaneous presence of three conditions is necessary: (1) susceptible material, (2) specific environment, and (3) tensile stresses. If one or more of these conditions is removed, EAC will not take place. Of the three main corrosion processes listed in Figure 36.1, the least

troubling is general corrosion since the propagation rate by uniform thinning of the container would be low and may not be life limiting for the containers. Localized corrosion and EAC may be more detrimental since these processes would tend to perforate the container at a faster rate at discrete locations allowing the ingress of water and the spreading of the radioactive material without substantial corrosion of the overall container wall. The occurrence of localized corrosion and EAC may be minimized by alloy selection, design, and fabrication.

E. REDUCING OR ANOXIC ENVIRONMENTS The containers in reducing environments will generally be surrounded by a backfill of bentonite, which will greatly limit the availability of oxygen to the metal surface. The lack of oxygen (or other oxidizing species) will create a redox potential that will be closer to the hydrogen evolution reaction. Elements such as iron (Fe), nickel (Ni), and copper (Cu) are mostly in the range of corrosion immunity at these reducing potentials in the near-neutral pH range [3, 30]. The most common materials under study in typically reducing environments are carbon steel, copper, and titanium [2, 3]. For the least corrosive underground waters, carbon steels could be viable materials; however, for the most saline conditions, titanium alloys are also being studied. E1. Carbon Steel and Low-Alloy Steel Carbon steels (and to a lesser extent low-alloy steels) have been tested in several countries in groundwater environments

REDUCING OR ANOXIC ENVIRONMENTS

507

Environmental or External Factors

Metallurgical or Internal Factors

(1) Composition of the electrolyte (2) Temperature (3) Redox potential (4) Effect of radiation, microbial activity and temperature on electrolyte composition, etc.

(1) Cast, welded or wrought metal (2) Processing factors (3) Phase transformation due to thermal aging, etc.

Affect the types of corrosion processes

If Ecorr < Ecrit

General Corrosion

If susceptible microstructure + specific environment + tensile stresses

If Ecorr > Ecrit

Localized Corrosion

Environmentally Assisted Cracking

Impact on the lifetime of the

Waste Package Container

FIGURE 36.1. Schematic representation of the effect of material characteristics plus type of environment on the lifetime performance of the waste package.

for the last 40 years. Carbon steels are rather inexpensive and they tend to suffer general corrosion rather than localized corrosion such as the more expensive stainless steels [31]. Most of the studies find that the corrosion rate of carbon steels and low-alloy steels is low, especially in anoxic conditions. Corrosion rates measured for carbon steels in granitic waters ranged from 1 to 55 mm/year, with one study showing that the rate reached a maximum at around 80 C [32]. The conditions that would lead to localized corrosion of carbon steels are quite specific and unlikely to be present in typical granitic groundwaters [33]. It was proposed that hydrogen embrittlement and hydrogen blistering of carbon steels may be possible if a high rate of hydrogen production exists [33]. However, recent calculations suggested that the failure of buried carbon steel containers by hydrogen damage mechanisms is unlikely [34]. The corrosion behavior of carbon steel was also studied in basaltic water, and it was found that even in oxygenated solutions at 150 C the corrosion rate of all tested carbon steels in basaltic waters was only on the order of 100 mm/year [35]. Under the Swedish program, researchers have studied the anoxic corrosion behavior of carbon steel and cast iron in groundwater at 50 and 85 C and the impact of the

presence of copper on the type and mechanical properties of the films formed on the iron alloys [36]. They used a barometric cell filled with a simulated groundwater and monitored the redox potential in the cell at 30 C on a gold electrode. Smart et al. determined that when steel was introduced to the cell, the redox potential decreased rapidly due to the consumption of the residual oxygen by the corrosion of the steel [36]. As part of the Japanese program of nuclear waste disposal, the passive corrosion behavior of steels was found to be dependent on variables such as groundwater pH, temperature, and available dissolved oxygen [37]. Fujiwara et al. have raised the concern that, whenever the corrosion of steel decreases due to a decrease in the oxygen content, the alkalinity in the immediacy of the steel increases. Since higher alkalinity would reduce the free corrosion potential of the steel, the process may increase the rate of hydrogen gas production, which could be detrimental for the stability of the repository [38]. Dong et al. have reported that the corrosion rate of carbon steel is dependent on the amount of bicarbonate (HCO3 ) present in the water [39]. At bicarbonate levels of 0.1 M, similar to the geologic disposal site, the corrosion of carbon steel is inhibited [39].

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Carbon steel has been identified as a candidate material for rock salt repositories in the German program. Studies were conducted to determine the corrosion response of welded and nonwelded Fe–1.5 Mn–0.5 Si steel in a MgCl2-rich brine (Qbrine) at 150 C under an irradiation field [40]. Welding was carried out by gas tungsten arc (GTAW) and electron beam (EB). The overall corrosion rate of both welded and nonwelded materials was approximately 70 mm/year; however, the welded materials experienced some localized attack in the weld seam area. When the material was heat treated for 2 h at 600 C, the corrosion rate of the welded material increased by approximately 40% [40]. Carbon steel and low-alloy steel have also been identified as candidate materials to contain nuclear waste for an intermediate storage of 100 years in the French program [41]. The dry oxidation testing of carbon steel in dry air (less than 15 ppm water), in air plus 2% water, and in air plus 12% water at 300 C for up to 700 h showed little damage to the tested coupons. When the depth of the oxide layer was extrapolated to 100 years, it resulted in less than 150 mm of damage. The authors also noted little or no water vapor effect on the oxidation rate at 300 C [41]. For the French repository, the container will be a cylinder 60 cm in diameter and 1.6 m long of carbon steel with an external wall thickness of 55 mm [42, 43]. E2. Copper The container for the disposition of nuclear waste in Sweden will consist of a 50-mm-thick layer of copper over cast nodular iron, which will provide the mechanical strength. Groundwater in granitic rock (as in the Swedish repository) is oxygen free and reducing below a depth of 200 m. The redox potential is between 200 and 300 mV on the hydrogen scale and the pH ranges from 7 to 9 [8, 44]. The chloride concentration in the groundwater can vary from 0.15 mM to 1.5 M with an equivalent amount of sodium and less calcium. The corrosion of a copper container in this reducing environment is expected to be less than 5 mm in 100,000 years of emplacement [44]. The corrosion of copper is mainly controlled by the availability of oxygen (trapped initially in the pores of the bentonite-based sealing materials), sulfate, and sulfide in the groundwater. The time to failure of the copper layer in the Swedish container has been modeled, and it is predicted that this failure, both by general and pitting corrosion, would be higher than 106 years under realistic emplacement conditions [45]. The anodic behavior of copper was also studied as part of the Japanese nuclear waste disposal program using potentiodynamic polarization tests in simulated groundwater at 30 C [46]. The amount of dissolved oxygen as well as different additions of chloride, sulfate, and bicarbonate was controlled. Imai et al. concluded that both sulfate and chloride promote the active dissolution of copper while carbonate is a passivating agent [46].

In the Canadian design, the thickness of the external copper layer is 25 mm. Scientists have modeled the failure mechanism of copper as a function of the oxygen availability, the temperature, the salinity of the solution, and the redox potential [7]. It is predicted that copper will undergo general corrosion and pitting during the initial warm and oxidizing period but only general corrosion during the subsequent longer anoxic cooler period. It has been predicted by this model that the Canadian copper container could last more than 106 years [47]. When a copper container is buried in a mostly reducing environment, the metal will initially be in contact with oxygen, until the oxygen is fully consumed, for example, by corrosion [48–50]. When all the oxygen is consumed, its corrosion potential will be in the region of immunity, that is, the thermodynamic oxidation of copper will not be possible. One of the few corrosion concerns about the use of copper in the repositories in Finland, Sweden, and Canada is that copper may be susceptible to EAC in waters containing, for example, ammonia and nitrite (NO2 ). This has been confirmed in studies for the Canadian program using the slow-strain-rate technique [51]. It has been reported that the crack growth rate could be as high as 8 nm/s [52]. However, the conditions under which the damage occurred were extreme and unrepresentative of container emplacement conditions. In the actual container, the general absence of aggressive SCC species, the limited applied strain, and the limited supply of oxygen will limit the susceptibility to environmental cracking. In another study, it has been shown that the minimum stress intensity for crack propagation in pffiffiffiffi copper for the Swedish container was 30 MPa m when tested inpaffiffiffiffi0.3 M NaNO2 solution [53]. A stress intensity of 30 MPa m can be considered high for a statically loaded container that may have shallow defects on the surface. That is, the conditions at which copper was cracked in the laboratory are too extreme to be representative of the actual repository conditions. E3. Stainless Steel and Nickel Alloys The cyclic potentiodynamic polarization method (ASTM G 61) was used to evaluate the anodic behavior of corrosionresistant alloys in oxidized Boom clay water (for the repository in Belgium) with varying degrees of added chloride at 90 C [54]. The original Boom clay water is dominated by chloride and sulfate. The alloys studied included 316L SS (also with high Mo and with Ti) (S31603), alloy 926 (N08926), alloy 904L (N08904), alloy C-4 (N06455), and Ti Gr 7 (R52400) (Table 36.2). It was found that both R52400 and N06455 resisted pitting corrosion even at added chloride concentrations of 10,000 ppm and N08926 resisted pitting up to 1000 ppm chloride. The other alloys showed minor pitting at 100 ppm chloride and definite pitting corrosion at the higher tested chloride concentrations [54].

OXIDIZING ENVIRONMENTS

E4. Titanium Titanium (Ti) alloys have been studied as candidate materials for the containers in Canada, Japan, and Germany. The titanium alloys were selected as a potential alternative because of their excellent performance in more aggressive brine solutions compared, for example, to stainless steels. The corrosion rates for Ti Gr 2 and Ti Gr 12 in both oxygenated and irradiated basalt environments are very low—less than 2 mm/year (0.08 mil/year) [55]. Shoesmith et al. also explained the failure mechanism and a predictive model for the degradation of Ti Gr 2 under the Canadian repository conditions [55]. The model takes into account the crevice propagation rate as a function of temperature and oxygen availability as well as other factors such as the amount of hydrogen absorbed by the alloy during corrosion before a critical concentration for failure is reached. The localized corrosion resistance of titanium alloys has also been investigated extensively as part of the Japanese program [56, 57]. Testing showed that as the temperature and the chloride concentration increased, the repassivation potential (ER, CREV) for Ti Gr 1 and Ti Gr 12 decreased to values well below the corrosion potential (Ecorr) [56]. Ti Gr 12 was more resistant to crevice corrosion than Ti Gr 1. For the other tested alloys, at constant temperature and chloride concentration, ER,CREV increased as the palladium (Pd) content in the alloy increased, rapidly up to 0.008% Pd and then slower between 0.008 and 0.062% Pd [57]. Titanium alloys were also investigated for their resistance to EAC. One way by which titanium alloys may suffer EAC under reducing conditions is by the formation of hydrides due to the slow absorption of hydrogen from the environment. Slow strain rate testing was conducted using Ti Gr 1 in deaerated 20% NaCl at 90 C at an applied potential of 1.2 V [standard hydrogen electrode (SHE)] [58]. It was confirmed that cracks initiated as deep as the presence of hydrides, that is, the presence of hydride was necessary for cracks to initiate. Based on the critical cracking thickness and the predicted amount of hydrogen generated, the authors dismissed the hypothesis that the titanium-made containers may fail by cathodic EAC [58].

F. OXIDIZING ENVIRONMENTS The case of the containers for the repository of Yucca Mountain is discussed separately from the other repositories since Yucca Mountain at this time is the only proposed repository that may have an unsaturated environment with unrestricted availability of air (oxygen). The design of the waste package for the Yucca Mountain repository has evolved in the last 15 years [2, 9]. In previous versions of the container, a thick layer of carbon steel was specified for the outer shell of the container and a corrosion-resistant material as the inner shell. However, since 1998, the design

509

of the engineered barriers has not changed significantly, and it currently specifies a double-walled cylindrical container covered by a titanium alloy drip shield. The outer shell of the container will be a Ni–Cr–Mo alloy (N06022) (Table 36.2), with an inner shell of nuclear-grade austenitic type 316 stainless steel (S31600). The function of the outer barrier is to resist corrosion and the function of the inner barrier is to provide mechanical strength and a shield to radiation. The drip shield will be made of Ti Gr 7 and a higher strength Ti alloy (Ti Gr 29) will be used for the internal ribs of the shields. The function of the drip shield is to deflect rock fall and early water seepage on the container [9]. F1. Corrosion Behavior of Alloy 22 (N06022) The container may suffer corrosion only if water is present in sufficient amount at the repository site. Dry corrosion of alloy 22 is negligible for the emplacement conditions. There are three main modes of corrosion that the container may suffer during its emplacement time (Fig. 36.1): (1) uniform, general, or passive corrosion, (2) localized corrosion (e.g., crevice corrosion), and (3) environmentally assisted cracking (e.g., SCC) [59]. All three types of corrosion may be influenced by the environment, including temperature, solution composition (chloride and nitrate concentration), redox potential, and the presence of microorganisms. F2. Uniform and Passive Corrosion of Alloy 22 General corrosion (or passive corrosion) is the uniform thinning of the container alloy at its open-circuit potential or corrosion potential (Ecorr). In the presence of aerated multi-ionic brines, such as those that may be present at the repository site, alloy 22 is expected to remain passive at its Ecorr. The passive corrosion rates of alloy 22 after 5 years immersion in multi-ionic solutions simulating concentrated groundwaters from pH 2.8 to 10 are extremely low and on the order of 10 nm/year [60–62]. This low corrosion rate was measured at 60 and 90 C for welded and nonwelded alloy 22 at an Ecorr range from 100 to þ 400 mV SSC (saturated silver chloride) electrode. The low corrosion rates or passive behavior of alloy 22 is because of the formation of a protective inner chromium-rich oxide film between the alloy (metal) and the surrounding electrolyte. This passive film is even stable in the presence of strong mineral acids at temperatures below 60 C [63]. It has been shown that the thickness of this passive film formed in concentrated hot electrolyte solutions could be only in the range of 5–6 nm [64]. The long-term extrapolation of the corrosion rate of alloy 22 has been modeled considering that the dissolution rate is controlled by the injection of oxygen vacancies at the oxide film–solution interface [65]. It has been concluded that it is unlikely that catastrophic failure of

510


FIGURE 36.2. Effect of temperature and chloride concentration on the crevice repassivation potential of alloy 22 using THE method.

the container may occur due to long-term passive film dissolution [65]. That is, the passive dissolution of alloy 22 is not considered to be the limiting factor for the life performance of the waste container. F3. Localized Corrosion of Alloy 22 Localized corrosion may be the most detrimental of the degradation modes in Figure 36.1. Localized corrosion (e.g., crevice corrosion) is a type of corrosion in which the attack progresses at discrete sites or in a nonuniform manner. The degradation model assumes that localized corrosion will only occur when Ecorr is equal or greater than a critical potential (Ecrit) for localized corrosion [66]. That is, if Ecorr < Ecrit, only general or passive corrosion will occur. Here, Ecrit can be defined as a certain potential above which the current density or corrosion rate of alloy 22 increases significantly and irreversibly above the general corrosion rate of the passive metal. The margin of safety against localized corrosion will be given by the value of DE ¼ Ecrit Ecorr. The higher the value of DE, the larger the margin of safety for localized corrosion. It is important to note here that the values of both Ecorr and Ecrit may depend both on the metallurgical condition of the alloy and the environment, such as temperature, chloride concentration, and the presence of inhibitors. The value of Ecorr is determined by measuring the long-term steady-state value of the open-circuit potential in each environment of relevance and Ecrit is the crevice repassivation potential measured using electrochemical techniques such as the cyclic potentiodynamic polarization (ASTM G 61) [66].

Localized corrosion was the most extensively studied mechanism of degradation in alloy 22, mainly since 2002 [67, 68]. Alloy 22 is extremely resistant to pitting corrosion but may suffer crevice corrosion, especially in pure chloride solutions and at temperatures higher than 60 C [69, 70]. Figure 36.2 shows the effect of chloride concentration and temperature on the repassivation potential of alloy 22, obtained using the Tsujikawa–Hisamatsu electrochemical (THE) method (ASTM G 192). The higher the temperature and chloride concentration, the lower the resistance of the alloy to localized corrosion [70]. The crevice corrosion susceptibility of alloy 22 is only promoted by the presence of chloride ions, and it can be fully inhibited by the presence of other anions in solution such as nitrate, bicarbonate, sulfate, fluoride, phosphate, and so on [71–73]. The best inhibitor of crevice corrosion in alloy 22 is nitrate [69, 74, 75]. The presence of inhibitor is generally stated using the ratio R ¼ [inhibitor]/[Cl ]. The higher is the value of R, the higher is the inhibition effect. Figure 36.3 shows the cyclic potentiodynamic polarization of alloy 22 in three different electrolyte solutions at 110 C. The electrolytes are (1) pure 8 m chloride (4 m NaCl þ 4 m KCl), (R ¼ 0) and (2) 8 m chloride with added sodium and potassium nitrate to obtain R ¼ 0.1 and R ¼ 0.5. Figure 36.3 shows that as R increased, the repassivation potential increased. That is, for R ¼ 0, Ecrit ¼ 210 mV SSC; for R ¼ 0.1, Ecrit ¼ 50 mV SSC; and for R ¼ 0.5, Ecrit ¼ 337 mV SSC. For R ¼ 0.5 there was a total inhibition of crevice corrosion (no hysteresis in the reverse scan in Figure 36.3). Since the groundwater at the repository site in Yucca Mountain contains not only chloride but also a variety of anions that may act as inhibitors for crevice

FIGURE 36.3. Effect of nitrate (ratio R ¼ 0, 0.1, 0.5) on repassivation potential of alloy 22 in 8 m chloride solution at 110 C.

OXIDIZING ENVIRONMENTS

FIGURE 36.4. Crevice corrosion under a crevice former in alloy 22 under constant applied potential of þ 100 mV SSC in 3.5 m NaCl þ 0.175 m KNO3 solution (R ¼ 0.05) at 100 C. Top part of image, interdendritic attack in weld metal; lower part, intergranular attack in base metal.

corrosion, it is unlikely that alloy 22 would suffer crevice corrosion under natural emplacement conditions [76]. Constant potential laboratory tests have shown that crevice corrosion in alloy 22 often repassivates after initiation due to a stifling mechanism [77]. Figure 36.4 shows the crevice corroded area under one crevice former tooth (ASTM G 192) after a week-long test in 3.5 m NaCl þ 0.175 m KNO3 solution (R ¼ 0.05) at 100 C at a constant applied potential of þ 100 mV SSC. The test was conducted at approximately 200 mV higher than the crevice repassivation potential of alloy 22 ( 110 mV SSC) in the same conditions [78]. Current measurements showed that crevice corrosion nucleated at 10 min after the potential was applied and it progressed with increasing anodic currents for the next 14 h, after which the anodic current started to decrease, becoming cathodic at hour 79 even though the potential was maintained at þ 100 mV SSC during the entire test [78]. Figure 36.4 also shows that crevice corrosion occurred under the entire footprint of the crevice former and that the depth of attack was shallow and even for both the base metal (lower part of the image) and the weld metal (upper part of the image). In the weld metal the attack was interdendritic and in the base metal the attack was intergranular [78]. F4. Environmentally Assisted Cracking of Alloy 22 Wrought mill annealed (MA) alloy 22 is highly resistant to EAC in most environments, including acidic concentrated

511

FIGURE 36.5. Anodic polarization curves for alloy 22 in SCW solution at 60 and 90 C, where an anodic peak is observed in the middle on the passive range of potentials.

and hot chloride solutions. Welded and nonwelded U-bend specimens of alloy 22 and five other nickel-based alloys exposed for more than five years to multi-ionic solutions that represent concentrated groundwater of pH 2.8 to 10 at 60 and 90 C were free from EAC [79, 80]. Even though alloy 22 is resistant to EAC in concentrated hot chloride solutions, it may be susceptible under other severe environmental conditions. Slow strain rate tests were performed using MA alloy 22 specimens in simulated concentrated water (SCW) and other solutions as a function of the temperature and applied potential [81, 82]. SCW has a pH 8–10, and it is approximately 1000 times more concentrated than groundwater. Alloy 22 was found susceptible to EAC in hot SCW solutions and bicarbonate plus chloride solutions at anodic applied potentials approximately 300–400 mV more positive than Ecorr. The occurrence of EAC was related to the presence of an anodic peak in the polarization curve of the alloy in SCW environments. For example, at ambient temperatures, the peak is not present and EAC does not take place [81]. Currently, the origin of the anodic peak is being investigated [83]. Figure 36.5 shows the polarization curves for alloy 22 in SCW solution, showing the presence of the anodic peak at 60 and 90 C. It was demonstrated that the most aggressive species for EAC in SCW was bicarbonate but that the presence of chloride in the bicarbonate solution enhances the aggressiveness of the environment [82]. Figure 36.6 shows typical transgranular cracking in alloy 22 tested in the laboratory using the slow strain rate technique under anodic polarization. The conditions at which SCC was found

512


while Ti Gr 12 generally exhibited the highest corrosion rates. Titanium and Ti alloys may be susceptible to EAC, such as hydrogen embrittlement (HE). Embrittlement by hydrogen is a consequence of absorption of atomic hydrogen by the metal to form hydrides. This may happen in service when the Ti alloy is coupled to a more active metal in an acidic solution. A critical concentration of hydrogen in the metal may be needed for HE to occur. Stress corrosion cracking was reported in Ti Gr 7 specimens subjected to constant-load tests in a concentrated groundwater solution pH 10 at 105 C [88]. Results from up to five-year immersion testing at 60 and 90 C of U-bend specimens made of wrought and welded Ti Gr 7 and Ti Gr 16 alloys showed that these alloys were free from EAC in multi-ionic solutions that could be representative of concentrated groundwater [89]. Welded Ti Gr 12 U-bend specimens suffered EAC in SCW liquid at 90 C. Under the same conditions, nonwelded Ti Gr 12 was free from cracking [80, 89]. FIGURE 36.6. Transgranular EAC fracture surface in alloy 22 after a slow strain rate test at 1.67 10 6 s 1 at an applied potential of þ 400 mV SSC in SCW at 65 C.

in the laboratory in alloy 22 are unrealistic for the natural conditions of the emplacement site in Yucca Mountain. F5. Corrosion Behavior of Titanium Alloys Titanium grade 7 (Ti Gr 7 or R52400) was selected to fabricate the detached drip shield for the repository in Yucca Mountain [9]. Other Ti alloys of higher strength such as Ti Gr 29 may also be used for the structural parts of the drip shield. The presence of the drip shield would deflect early water seepage from the containers. This drip shield would also deflect rock fall from the containers. Ti Gr 7 belongs to a family of Ti alloys especially designed to withstand aggressive chemical environments (Table 36.2) [84]. The superior corrosion resistance of Ti and Ti alloys is due to a thin, stable, and tenacious oxide film that forms rapidly in air and water, especially under oxidizing conditions. A detailed review of the general, localized, and environmentally assisted cracking behavior of Ti Gr 7 and other titanium alloys relevant to the application in Yucca Mountain has addressed, among other topics, the effect of alloyed palladium, the properties of the passive films, and the effect of radiation [85]. The presence of fluoride in the groundwater may render Ti Gr 7 more susceptible to general and crevice corrosion under anodic polarization [86]. Weight loss, creviced, and U-bend specimens of Ti Gr 7, 12, and 16 were exposed to three different aerated electrolyte solutions simulating concentrated groundwater for over five years at both 60 and at 90 C in the vapor and liquid phases of these solutions [87]. Ti Gr 7 generally exhibited the lowest corrosion rates irrespective of temperature or solution type

G. SUMMARY AND CONCLUSIONS 1. The consensus around the world is that high-level nuclear waste should be deposited in stable geologic repositories and several countries are currently developing them. 2. Most of the repositories in the world are planned to be in stable rock formations (e.g., granite) below the water table (saturated). The United States is studying a repository above the water table (unsaturated) and with unrestricted access of oxygen. 3. The repositories will consist of a stable geologic formation within which engineered barriers will be constructed. The most important part of the engineered barrier is the container for the waste 4. The containers are, in general, designed as doublewalled metallic cylinders. Some carbon steel containers may be single walled. 5. From the corrosion point of view, most repositories will have reducing or anoxic environments. The U.S. repository will have a natural oxidizing environment. 6. Copper, titanium, and carbon steels were determined to be suitable materials for the reducing or anoxic repositories. 7. High-end materials such as alloy 22 and Ti Gr 7 are being characterized for the mostly dry and oxidizing environment of the U.S. repository. 8. Materials for the engineered barriers are being evaluated for general corrosion, localized corrosion, and environmentally assisted cracking resistance. General corrosion is not generally an important factor that determines the lifetime performance of the containers. Environmentally assisted cracking and localized corrosion are more detrimental and less predictable.

REFERENCES

9. The candidate materials for the containers are studied under different metallurgical conditions, such as mill annealed, welded, and thermally aged. 10. Some of the most important environmental variables that may affect the corrosion behavior of the engineering materials include the concentration and type of the aqueous electrolytes at the site, temperature, and redox potential established.

13.

14.

REFERENCES 1. International Atomic Energy Agency, www.iaea.org. 2. R. B. Rebak and R. D. McCright, “Corrosion of Containment Materials for Radioactive Waste Isolation,” in ASM Handbook, Vol. 13C, Corrosion: Environments and Industry, ASM International, Metals Park, OH, 2006, pp. 421–437. 3. D. W. Shoesmith, Corrosion, 62, 703 (2006). 4. D. G. Bennett, A. J. Hooper, S. Voinis, and H. Umeki, “The Role of the Engineered Barrier System in Safety Cases for Geological Radioactive Waste Repositories: A NEA Initiative in Co-operation with the EC,” in Proceedings from the XXIX Symposium on Scientific Basis for Nuclear Waste Management, Vol. 932, Materials Research Society, Warrendale, PA, 2006, pp. 43–52. 5. Nuclear Energy Agency, “Engineered Barrier Systems and the Safety of Deep Geological Repositories,” Organisation for Economic Co-Operation and Development OECD Publications, Paris, France, 2003. 6. P. A. Witherspoon and G. S. Bodvarsson, “Geological Challenges in Radioactive Waste Isolation—Third Worldwide Review,” Report LBNL-49767, Lawrence Berkeley Laboratory, Berkeley, CA, 2001. 7. D. W. Shoesmith, B. M. Ikeda, F. King, and S. Sunder, “Prediction of Long Term Behavior for Radioactive Nuclear Waste Disposal,” in Research Topical Symposia—Life Prediction of Structures Subject to Environmental Degradation, NACE International, Houston, TX, 1996, p. 101. 8. L. O. Werme, “Fabrication and Testing of Copper Canister for Long Term Isolation of Spent Nuclear Fuel,” Vol. 608, Materials Research Society, Warrendale, PA, 2000, p. 77. 9. G. M. Gordon, Corrosion, 58, 811 (2002). 10. R. L. Clark, “Proposed Amendments to the Environmental Radiation Protection Standards for Yucca Mountain, Nevada” in 11th International High Level Radioactive Waste Management Conference Proceedings, Las Vegas, NV, April 30–May 4, 2006, American Nuclear Society, La Grange Park, IL, 2006, pp. 1124–1130. 11. P. N. Swift, K. Knowles, J. McNeish, C. W. Hansen, R. L. Howard, R. MacKinnon, and S. D. Sevougian, “Long-Term Performance of the Proposed Yucca Mountain Repository, USA,” in Proceedings from the XXXII Symposium on Scientific Basis for Nuclear Waste Management, Vol. 1124, Materials Research Society, Warrendale, PA, 2009, pp. 3–14. 12. A. A. Sag€ues, “Nuclear Waste Package Corrosion Behavior in the Proposed Yucca Mountain Repository,” in Scientific Basis

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37 CORROSION BEHAVIOR OF ELECTRODEPOSITED NANOCRYSTALS U. ERB Department of Materials Science and Engineering, University of Toronto, Toronto, Ontario, Canada

A. B. C. D. E.

Introduction Synthesis of nanomaterials by electrodeposition Structure of nanocrystalline electrodeposits Mechanical and physical properties of nanomaterials Corrosion properties of nanocrystalline electrodeposits E1. Nanocrystalline nickel and nickel-based alloys E2. Nanocrystalline cobalt and cobalt alloys E3. Nanocrystalline zinc E4. Nanocrystalline copper F. Industrial applications and outlook References

A. INTRODUCTION This chapter deals with the corrosion behavior of nanocrystalline materials made by electrodeposition. The synthesis method, structure of nanodeposits, and some mechanical and physical properties of these materials will be briefly described. This will be followed by a review of advances made over the past 20 years in the understanding of the corrosion properties of nanocrystalline nickel, cobalt, zinc, and copper and some of their alloys. It will be shown that, contrary to earlier concerns, the high concentrations of intercrystalline defects in nanocrystalline metals do not compromise their corrosion resistance. Over the past 30 years, numerous synthesis methods have been developed to make nanocrystalline materials in different shapes and forms. From a processing point of view, five basic approaches are used to achieve microstructural refinement down to the 1–100 nm range: vapor-phase processing,

liquid-phase processing, solid-state processing, chemical synthesis, and electrochemical synthesis. For each approach there are a variety of individual methods that can make very different or similar materials (Table 37.1). In terms of microstructures, nanomaterials can be classified according to their dimensional structure modulation. In one particular structure scheme [1], zero-dimensional nanomaterials are individual clusters, particles, or fibres with any aspect ratio. One-dimensional nanomaterials are layered (often epitaxially grown) structures with layer thicknesses less than 100 nm. Twodimensional nanomaterials are thin layers with a grain size less than 100 nm, while three-dimensional nanomaterials are bulk materials with average grain size less than 100 nm. For bulk three-dimensional nanomaterials there are four basic structure types, depending on the synthesis approach. For materials that are made by the consolidation of precursor nanoparticles (e.g., produced by inert gas condensation or ball milling), the main microstructural defects are grain boundaries (between the consolidated particles) and residual porosity (interparticle voids). In materials made by crystallization from amorphous precursors there are typically several crystalline phases plus, in many cases, a residual amorphous phase separating individual crystals. Nanomaterials made by methods such as severe plastic deformation contain mainly grain boundaries and high densities of dislocations. Finally, in nanomaterials made by techniques such as electrodeposition, the main structural defects are grain boundaries and triple junctions (the lines between three adjoining crystals). In view of the large number of different synthesis techniques for nanomaterials, their complexity in terms of dimensional structure modulations and the various structure types even within one single group of bulk nanomaterials, it is perhaps not surprising that a comprehensive treatise on the corrosion properties of these exciting new materials is


517

518

CORROSION BEHAVIOR OF ELECTRODEPOSITED NANOCRYSTALS

TABLE 37.1. Five Major Processing Routes for Making Nanomaterials with Specific Examples Processing Route

Specific Examples

Vapor-phase processing

Physical vapour deposition Chemical vapour deposition Inert gas condensation Rapid solidification Atomization Sonication of immiscible liquids Annealing of amorphous materials Mechanical attrition Equal channel angular processing Sol–gel processing Precipitation Inverse micelle technology Electrodeposition Galvanic displacement Electroless plating

Liquid–phase processing

Solid–state processing

Chemical synthesis

Electrochemical synthesis

currently not available. While some earlier studies of sputterdeposited nanocrystalline 304 stainless steel–type Ni–Fe–Cr films showed very promising results in improving the corrosion resistance in 0.3 wt % NaCl solution compared with their conventional polycrystalline counterparts [2], another study using magnetron-sputtered nanocrystalline Fe–8 wt % Al showed both detrimental and beneficial effects of microstructural refinement with respect to corrosion resistance depending on the environment [3]. Two of the most intensively studied subgroups of nanomaterials in terms of their corrosion properties are materials made by (i) crystallization of amorphous precursors and (ii) electrodeposition methods. In the first group of nanocrystallized materials (see e.g., [4–9] for some of the early studies) the interpretation of the electrochemical response of various materials in different environments is often rather difficult because the nanostructured materials may contain several crystallographically and chemically different phases in addition to crystal size refinement and residual amorphous phase. On the other hand, materials made by electrodeposition can be produced both in conventional polycrystalline and nanocrystalline forms without changes in phase composition and without residual amorphous phase. Therefore, the assessment of microstructural refinement down to the nanoscale (i.e., crystal size effects) on the corrosion performance is much easier in this group of materials. This chapter summarizes our current understanding of the corrosion behavior of nanocrystalline materials made by electrodeposition. Of particular interest will be the following questions: 1. To what extent does grain size reduction influence the electrochemical response of certain groups of pure metals and alloys?

2. What is the effect of grain size on the corrosion morphology? 3. What is the overall gain/loss of corrosion resistance of these materials when we go from conventional polycrystalline structure to nanostructures? The chapter is organized as follows. First, the synthesis of nanocrystalline materials by electrodeposition will be briefly reviewed. Next, we will look at the unique microstructural features of these materials in comparison with their conventional polycrystalline counterparts. After a brief review of their mechanical and physical properties, the main focus of this chapter will be the corrosion behavior of nanocrystalline nickel, cobalt, copper, and zinc electrodeposits as well as some of their alloys and composites. The last section of this chapter will address some of the applications of these materials and the future outlook for this group of nanomaterials.

B. SYNTHESIS OF NANOMATERIALS BY ELECTRODEPOSITION Electrodeposition is one of several methods in the general group of electrochemical synthesis of nanomaterials. Other processes using the electrochemical approach include electroless deposition, galvanic displacement, or electrodeposition under oxidizing conditions. In this chapter we will limit our discussion to electrodeposition of metals and alloys from aqueous solutions. Molten salt electrolysis, ionic liquid deposition, or plating of ceramics and semiconductors will not be considered here. The potential for making metals, alloys, and composites in industrially and economically viable electroplating operations has been established throughout the 1980s and 1990s. The early work focused on nickel and nickel-based alloys. This was followed by expanding the technology to cobalt, copper, and zinc and several of their alloys (see Table 37.2 for examples). The first patents on this technology were issued in 1994 and 1995 [10, 11] and the first large-scale industrial application of an electrodeposited nanocrystalline nickel–phosphorus alloy was reported in 1996 [12]. For a more comprehensive treatment on the synthesis of electrodeposited nanomaterials the reader is referred to recent review articles published on this subject [13–16]. Electrodeposition of a pure metal from aqueous solutions at low pH containing Mez þ ions involves the following cathodic reactions: Mez þ þ ze ! Me H þ þ e ! 12 H2

ðcathodic reactionÞ ðcathodic reactionÞ

ð37:1Þ ð37:2Þ

SYNTHESIS OF NANOMATERIALS BY ELECTRODEPOSITION

TABLE 37.2. Examples of Electrodeposited Materials with Nanocrystalline Structures Ni Ni–P Ni–Fe Ni–Zn Ni–Fe–Cr Ni–Zn–P Ni–SiC Ni–Al2O3 Ni–P–BN Ni–MoS2 Ni–Al (particles) Ni–Carbon nanotubes

Co Co–P Co–W Co–Fe Co–Fe–P Cu Cu–Al2O3 Zn Zn–Ni Fe–Co–Ni Pd Pd–Fe

Alloys can be made by adding two or more species to the plating bath. For example, for Zn–Ni alloys, the bath contains Zn2 þ and Ni2 þ ions which come from appropriate metal salts (e.g., sulfates, chlorides) dissolved in the electrolyte in various concentrations, depending on the required composition of the final alloy. In order to maintain the required metal ion solution in the plating bath, two approaches are commonly used. When a dimensionally stable anode (DSA) is used, the bath must be periodically replenished with metal salt additions. On the other hand, when the metal to be deposited is also used as the anode, continuous anode dissolution replenishes the bath with metal ions: Me ! Mez þ þ ze

ðanodic reationÞ

ð37:3Þ

Cathodic and anodic reactions can be more complex than shown in Eqs. 37.1–37.3 depending on the metal. Several intermediate steps could be involved, but this is beyond the scope of this chapter. Electrodeposition can also be used to make composite materials consisting of a metal matrix with a second-phase material embedded in the matrix in various amounts. This is usually achieved by adding the second phase (e.g., particles of SiC, Al2O3 or BN, or carbon nanotubes) to the plating bath, from which it is then codeposited with the metal. Electrodeposition involves various steps, including diffusion of ions from the bulk of the electrolyte to the cathode and through the Nernst diffusion layer, formation of first adions on the cathode surface, surface diffusion of adions, nucleation of crystals, and growth of crystals. Initially the structure of the cathode is very important (e.g., surface cleanliness, presence of oxides, surface step structures, dislocations, grain boundaries), but with increasing deposit thickness the cathode surface influence diminishes. Nanocrystalline metals require deposition under conditions that favor nucleation of new crystals and reduce growth

519

of existing crystals. Through several experimental and theoretical studies [17–20] it has been shown that massive nucleation can be maintained during the growth of the deposit under the conditions of high overpotential (i.e., current density during the plating process) and reduced adion mobility. The latter can be controlled within certain limits by the addition of surface-active elements (e.g., saccharin, coumarin, for the case of nickel plating) to the plating bath. However, it must be noted that the breakdown of such additives can result in the incorporation of impurities in the deposit. For example, the use of saccharin as a grain refiner in nickel deposition usually results in sulfur and carbon impurities in the deposit with concentrations on the order of several hundred parts per million [17]. Such impurities can have a significant effect on the corrosion behavior of electrodeposited nanomaterials. In electrodeposition with conventional direct current, the limiting current density of the system may not allow the application of high enough current densities to induce the massive nucleation required for nanocrystal formation. The limiting current density is reached when the metal ions are plated out at a rate higher than the rate with which they arrive in the Nernst diffusion layer by diffusion from the bulk of the electrolyte. In such cases, cathodic reaction 37.2 dominates and the quality of the electrodeposit deteriorates rapidly with increasing current density. To overcome this problem, two approaches are usually used. In the first approach, extensive agitation of the electrolyte can enhance the replenishment of metal ions in the Nernst diffusion layer. In the second approach, a very high current density is applied for only a short period of time. This is followed by a current-off period during which no plating takes place but which allows metal ions to diffuse back from the bulk of the electrolyte to the Nernst diffusion layer. This approach requires the use of a pulsed current power supply. Pulse plating of nanomaterials is typically carried out with current on and off times on the order of microseconds to milliseconds [10, 17, 18]. Electrodeposition is mainly used to apply a coating on a finished product to enhance certain properties of a part such as wear resistance, corrosion resistance, oxidation resistance, optical appearance, or magnetic properties. For such applications the thicknesses of the coatings are on the order of 0.01–1.0 mm depending on the application. Many industries use electrodeposited coatings for parts in automotive, aerospace, power generation, defense, and consumer products. The same infrastructure that already exists for conventional electrodeposition processes can also be used to make nanocrystalline electrodeposits. However, electrodeposition is not limited to nanocrystalline coatings. Special processes have been developed to make nanocrystalline thin foils, meshes, wires, thick structural plates, and components for microelectromechanical systems [21–23].

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C. STRUCTURE OF NANOCRYSTALLINE ELECTRODEPOSITS Schematic cross sections showing the microstructures of conventional polycrystalline and nanocrystalline electrodeposits are shown in Figure 37.1. When electrodeposition is carried out at relatively low current densities using directcurrent plating, the initial layer close to the substrate contains numerous small crystals. However, with increasing thickness the anisotropy in crystal growth results in a structural transition in which certain crystals grow rapidly while the growth of others is suppressed. Often this results in a columnar crosssectional structure with grain sizes in the micrometer range. On the other hand, nanocrystalline electrodeposits produced under conditions leading to massive nucleation throughout the entire plating process maintain the nanocrystalline grain structure over the entire thickness. Figure 37.2 shows both a planar view and a cross-sectional view of nanocrystalline nickel produced by pulsed current electrodeposition. The average grain size of this material is on the order of 15 nm. The structure is more or less equiaxed. The main structural defects in electrodeposited nanomaterials are grain boundaries, separating two crystals with different orientations, and triple junctions, the lines where three crystals meet. These defects are regions in the material with reduced structural order compared with the perfect crystal. Grain boundaries and triple junctions are typically 1 nm thick and vary in their structure depending on the orientations of the adjacent crystals. All polycrystalline materials contain a certain volume fraction of atoms associated with grain boundaries and triple junctions. Detailed calculations [24] have shown that for conventional polycrystalline materials with grain sizes larger than 1 mm, these volume fractions are very small (Table 37.3). On the other hand, for grain sizes less than 100 nm their volume fraction increases very rapidly, reaching close to 50% at a grain size of 5 nm. In other words, in a nanomaterial with 5 nm grain size half of the material is made up of atoms at grain boundaries and triple junctions.

FIGURE 37.1. Schematic cross sections showing microstructures of conventional polycrystalline (a) and nanocrystalline (b) electrodeposits. TABLE 37.3. Volume Fractions of Atoms Associated with Grain Boundaries and Triple Junctions as Function of Grain Size Assuming Grain Boundary Thickness of 1 nm Grain Size 100 mm 10 mm 1 mm

Volume Fraction 5

2.9 10 2.9 10 4 2.9 10 3

Grain Size 100 nm 10 nm 5 nm

Volume Fraction 0.003 0.271 0.488

In many polycrystalline materials grain boundaries are preferentially attacked during corrosion because of (i) their enhanced energy compared to the perfect crystal, (ii) the less than perfect atomic structure, and (iii) their susceptibility to segregation of impurities and formation of second-phase particles. For this reason, the initial expectation was that nanocrystalline metals made by electrodeposition would exhibit very poor corrosion properties mainly because of their high-grain-boundary-volume fractions. In Section E, it

FIGURE 37.2. Transmission electron micrographs of nanocrystalline nickel electrodeposit in planar (a) and cross-sectional view (b).

CORROSION PROPERTIES OF NANOCRYSTALLINE ELECTRODEPOSITS

TABLE 37.4.

521

Effect of Grain Size on Various Properties for Electrodeposited Nickel and Cobalt

Property

Poly Ni (Grain Size 10 mm)

Nano Ni (Grain Size 10 nm)

Poly Co (Grain Size 5 mm)

Nano Co (Grain Size 10 nm)

140 103 403 207 37 11 7.5 502

650 800 1100 204 21 10.5 13.5 488

232 311 805 207 40.6

525 1002 1865 200 37.0

7.5 1340

16.2 1380

Hardness (VHN) Yield strength (MPa) Tensile strength (MPa) Young’s modulus (GPa) Taber wear index Thermal expansion (10 6/K) Electrical resistivity [mW. cm] Saturation magnetization (kA/m)

will be shown that this is not the case for most electrodeposited nanomaterials studied to date.

D. MECHANICAL AND PHYSICAL PROPERTIES OF NANOMATERIALS The widespread interest in nanocrystalline materials is largely due to their outstanding mechanical properties and some of their unique property combinations which are not achievable in their conventional or amorphous counterparts. For nanocrystalline metals produced by electrodeposition it is found that grain size reduction to the nanometer range has no major effect on properties such as the Young’s modulus, thermal expansion, heat capacity, or saturation magnetization [14]. On the other hand, hardness, yield strength, ultimate tensile strength, wear resistance, and electrical resistivity are strongly affected by grain size [14]. Table 37.4 summarizes some of these properties for nanocrystalline nickel and cobalt in comparison with their polycrystalline counterparts. The substantial increases in their hardness and strength can be understood on the basis of the Hall–Petch effect [25, 26], which describes such increases in terms of grain boundary–dislocation interactions. Higher grain boundary densities hinder easy slip of dislocation, which makes materials harder and stronger. The Taber wear index is a measure of the material’s resistance to abrasive wear; the lower the number, the higher the wear resistance. It has been shown [27] that the Taber wear index of nanocrystalline nickel electrodeposits is directly related to their increased hardness. Directionally similar results as shown in Table 37.4 for nanocrystalline nickel and cobalt have also been observed for other electrodeposited metals and alloys such as Ni–P, Zn–Ni, Ni–Fe, and Cu.

E. CORROSION PROPERTIES OF NANOCRYSTALLINE ELECTRODEPOSITS When nanocrystalline metal electrodeposits were initially developed, their corrosion behavior was of great concern

because of their high density of grain boundaries and triple junctions. In conventional polycrystalline materials, these defects are often prone to intergranular and stress corrosion cracking mainly because of the energetic and chemical composition reasons discussed in Section C. This section reviews recent advances in the understanding of the corrosion properties of several pure nanocrystalline metals and some of their alloys. It will be shown that grain size reduction in many of these materials can actually improve their resistance to localized corrosion quite substantially. E1. Nanocrystalline Nickel and Nickel-Based Alloys In the early 1990s Rofagha et al. [28, 29] presented the first systematic study on the effect of grain size (100 mm, 500 nm, 50 nm, and 32 nm) on the corrosion behavior of nickel in deaerated 2 N H2SO4 solution at pH 0 using potentiodynamic and potentiostatic polarization tests. Figure 37.3 shows that nanocrystalline nickel exhibited the same active, passive, transpassive behavior as conventional polycrystalline nickel. However, the current density for the nanomaterials in the passive region was about one order of magnitude higher than for polycrystalline nickel. X-ray photoelectron spectroscopy showed that this higher passive current density was the result of a more defective passive film that formed on the nanocrystalline nickel surfaces [30]. Figure 37.3 also shows a positive shift in the corrosion potentials for all nanocrystalline materials which was explained in terms of the catalysis of the hydrogen evolution reaction. Despite the enhanced corrosion rate, Rofagha et al. [28, 29] showed that the corrosion morphology on the nanomaterials was very uniform instead of developing deep localized attack as observed along grain boundaries and triple junctions in polycrystalline nickel. Wang et al. [31] studied the corrosion behavior of polycrystalline Ni (50 mm) and nanocrystalline Ni (32 nm, 16 nm) in 30 wt % KOH (pH 14.8) at 24 C. Within a wide potential range from the hydrogen evolution reaction to the oxygen evolution reaction, all materials were very inert with low passivation currents. However, as for the tests in 2 N

522


1500 1250

Potential (mV)

1000 100 µm

32 nm

750 500 nm

50 nm

500 250 0 −250 −500 10−2

10−1

100

101

102

103

104

105

106

Current Density µA / sq.cm

FIGURE 37.3. Potentiodynamic polarization curves for polycrystalline (100 mm) and nanocrystalline (500 nm, 50 nm, and 32 nm) nickel in 2 N H2SO4 solution. (Reproduced from [29] with permission.)

H2SO4 [28, 29], the current densities for the nanocrystalline materials were somewhat enhanced, about 2–5 times higher than for polycrystalline nickel. The nanocrystalline materials again showed very uniform attack. Tang et al. [32] compared the corrosion performance of nanocrystalline nickel electrodeposits (grain sizes in the 5–10 nm range) made by different plating methods: pulse plating, direct-current plating, and pulse-reverse plating. Immersion tests were carried out in the following solutions: 7 M nitric acid, 3 M hydrochloric acid, and 20 g/dm3 citric

acid. In addition, samples were exposed to a moist SO2 environment. Considerable differences in the corrosion performance were observed and the results were interpreted in terms of distinct changes in the crystallographic textures in the different materials. Nanocrystalline nickel also showed excellent corrosion resistance when exposed to a salt spray environment as per American Society for Testing and Materials (ASTM) B117 [33]. In this study both nanocrystalline (grain size 10 nm) and polycrystalline (grain size 10 mm) nickel electrodeposits were prepared as 10-mm-thick coatings on mild steel substrates. Percentages of areas covered with red rust were recorded for exposure times up to 250 h. Both the nanocrystalline and the polycrystalline coatings provided the same protection against corrosion of the steel substrate, with area percentages of red rust approaching 30% after 250 h. In a more recent study, Kim et al. [34] investigated the effects of grain size and sulfur solute segregation on the corrosion behavior of polycrystalline (grains size 100 mm) and nanocrystalline nickel (grain size 20–30 nm), both containing about 1000 ppm by weight of sulfur impurities. Corrosion tests were carried out in 0.25 M Na2SO4 solutions at a pH of 6.5. Both materials showed very similar electrochemical behavior in potentiodynamic polarization curves. However, considerable differences were observed in their corrosion morphologies. Figure 37.4 shows that nanocrystalline nickel deposits developed a morphology consisting of numerous shallow corrosion pits ( 1 mm) and a nanocrystalline (grain size 100 nm) Ni–SiC nanocomposite coatings using sliding-type wear testing and electrochemical impedance spectroscopy in 0.5 M Na2SO4 neutral solution. They showed that the nanocrystalline nickel composite exhibited a higher polarization resistance and a 50% reduced corrosion rate, already in the absence of wear action. Even higher differences in material removal rates were found when the material was subjected to wear/corrosion conditions. For example, at a load of 30 N the nanocomposite materials showed a 90% reduction in material removal rate as compared with polycrystalline nickel. Peng et al. [42] produced nanocrystalline nickel electrodeposits (grain sizes 42, 38, 31 nm) with varying amounts (0, 4.5, 10.9 wt %) of chromium nanoparticles (average particle size 39 nm) by a codeposition process. Potentiodynamic polarization tests in 3.5% NaCl solution showed that the chromium particles reduced the corrosion potential, increased the breakdown potential, and basically eliminated pitting at the highest chromium concentration. X-ray photoelectron spectroscopy showed that at 10.9 wt % chromium a continuous passive chromium oxide film was formed on the surface. E2. Nanocrystalline Cobalt and Cobalt Alloys Kim et al. [43] studied the corrosion behavior of electrodeposited cobalt nanodeposits (grain size 13 nm) in 0.25 M Na2SO4 (pH 7) and compared the results with conventional cobalt (grain size 10 mm). Neither the polycrystalline nor the

524

CORROSION BEHAVIOR OF ELECTRODEPOSITED NANOCRYSTALS 0.8 0.6

Potential (VSCE)

0.4 0.2 0.0 -0.2

Annealed Poly-Co (~10 µm) Nano-Co (13 nm)

-0.4 -0.6 -0.8 -1.0 -1.2 10-7

10-6

10-5

Current Density

10-4

10-3

10-2

(A/cm2)

FIGURE 37.5. Potentiodynamic polarization curves in 0.1 M NaOH (pH 13) for polycrystalline and nanocrystalline cobalt.

nanocrystalline cobalt showed passivation in this solution and the potentiodynamic curves for both grain sizes were nearly identical. However, as observed before for the corrosion behavior of nickel, the nanocrystalline cobalt showed high resistance to localized attack. The same two cobalt materials were also tested in a 0.1 M NaOH solution (pH 13) in which conventional polycrystalline cobalt readily shows passive film formation [44]. Figure 37.5 presents the potentiodynamic polarization curves for both materials in this solution. It can be readily seen that the polarization curve for the nanocrystalline cobalt is very similar to the one for polycrystalline cobalt. In other words, grain size reduction from 10 mm to 13 nm did not substantially change the passivation behavior of cobalt. Aledresse and Alfantazi [45] also compared the corrosion behavior of polycrystalline (grain size 100 mm) and nanocrystalline (grain size 67 nm) cobalt in 0.25 M Na2SO4 solution. In addition, their study also included a nanocrystalline Co–P alloy with a grain size of 50 nm. None of the materials passivated in this solution. Jung and Alfantazi [46] performed electrochemical impedance spectroscopy as well as potentiostatic and potentiodynamic polarization tests on microcrystalline cobalt and nanocrystalline cobalt (grain size 20 nm) and cobalt–1.1 wt % phosphorus (grain size 10 nm) in 0.1 M H2SO4 solution. Potentiodynamic polarization results showed that all materi-

als exhibited active dissolution with no transition to passivation. While the polarization curves for nanocrystalline and microcrystalline cobalt were almost identical, a noticeable shift to more positive potential was observed for the nanocrystalline cobalt–phosphorus alloy. In addition, the overpotential for hydrogen evolution was observed to decrease. X-ray photoelectron spectroscopy showed that the enhanced corrosion resistance of the Co–P alloy was due to an enrichment of P of the corroded surface after polarization. However, at higher anodic overpotentials the benefit of phosphorus disappeared due to the formation of a nonprotective surface film which contained elemental phosphorus, hypophosphite, and phosphate species. Saito et al. [47] studied the corrosion performance of nanocrystalline Co65Ni12Fe23 ternary alloys (grain sizes 10–40 nm) in deaerated 2.5 wt % NaCl solution. Potentiodynamic polarization curves showed that the material passivated in this solution with passivation current densities of 30 mA/cm2. E3. Nanocrystalline Zinc Youssef et al. [48] compared the corrosion behavior of nanocrystalline zinc electrodeposits (grain size 56 nm) and conventional electrogalvanized zinc (grain size 8–20 mm) in deaerated 0.5 N NaOH solution at 25 C. Both potentiodynamic polarization and alternating current (ac) impedance measurements were used. The polarization curves for both materials were very similar in shape but showed differences in specific electrochemical parameters such as corrosion potential (Ecorr), passivation potential (Ep), maximum current density (im) and passive current density (ip), as summarized in Table 37.5. Both icorr and ip were lower for nanocrystalline zinc than for polycrystalline electrogalvanized zinc. It was concluded that the oxide film on nanocrystalline zinc was more protective than the film on polycrystalline film which was supported by ac impedance measurements. The average capacitance value for nanocrystalline zinc was 69 mF/cm2, as compared to 227 mF/cm2 for the polycrystalline material. It was further shown that nanocrystalline zinc displayed numerous discrete corrosion pits after potentiodynamic polarization while polycrystalline zinc showed a more uniform corrosion morphology. The corrosion properties of a series of electrodeposited zinc–nickel alloy coatings on mild steel substrates were

TABLE 37.5. Important Electrochemical Parameters for Nanocrystalline and Polycrystalline Zinc from Polarization Curves in Deaerated 0.5 N NaOH Solution Material

Ecorr (mV)

Nano Zn Poly Zn

1470 1455

Source: From [48].

Icorr (mA/cm2) 90 229

Ep (mV)

im (mA/cm2)

ip (mA/cm2)

1362 1342

4503 3895

210 828

INDUSTRIAL APPLICATIONS AND OUTLOOK

525

1.0 0.8

PC ~ 1000nm UFG ~ 250nm NC ~ 45nm

Potenital (VSCE)

0.6 0.4 0.2 0.0 -0.2 -0.4 -0.8 -0.8 -1.0 -1.2 10-9

10-8

10-7

10-6

10-5

10-4

10-3

10-2

Current Density (A/cm2)

FIGURE 37.6. Effect of grain size in electrodeposited copper on potentiodynamic polarization curves in 0.1 M NaOH solution (pH 13).

studied by Alfantazi and Erb [49]. This study covered both polycrystalline and nanocrystalline materials. It was shown through salt spray testing (ASTM B 117-81) that nanocrystalline Zn-45 w % Ni (grain size 20 nm) and Zn-63 wt % Ni (grain size 2 nm) showed greater resistance to the formation of white and red rust than polycrystalline Zn. However, in this study it was difficult to separate grain size effects from contributions due to chemical composition. E4. Nanocrystalline Copper Tao and Li [50] studied mechanical and electrochemical properties of nanocrystalline copper electrodeposits (grain size 56 nm) in comparison with polycrystalline copper deposits (grain size 2 mm). Potentiodynamic polarization curves for both materials in 0.1 M NaOH solution showed two differences. First, the corrosion potential of nanocrystalline copper was shifted slightly to more negative values. Second, the current density in the passive region was lower for the nanocrystalline copper than for the polycrystalline copper. On the other hand, Yu et al. [51] observed that grain size reduction in copper electrodeposits from 1 mm to 45 nm in the same solution (0.1 M NaOH) had no significant effect on the potentiodynamic polarization of copper, as shown in Figure 37.6. After photodynamic polarization tests all materials were covered with a fine needlelike corrosion product (Fig. 37.7) with a morphology similar to what in previous studies on polycrystalline copper was described as the CuO/Cu(OH)2 upper layer [52, 53]. Yu et al. [54] investigated the polarization behavior of nanocrystalline and polycrystalline copper in 3.5 wt % NaCl solution. They observed similar polarization behavior for both materials but noted that the nanocrystalline material had a higher Ecorr and a lower anodic current density which they attributed to the large density of intercrystalline defects

FIGURE 37.7. SEM images after potentiodynamic polarization tests in 0.1 M NaOH solution of polycrystalline and nanocrystalline copper revealed similar “needle like” corrosion product morphologies. Grain sizes: (a) 45 nm, (b) 250 nm and (c) 1 mm.

which act as preferential attack sites and lead to the rapid formation of a passive film. Unfortunately, in their study Yu et al. [54] did not provide any microstructural details on the materials used for the corrosion study. In fact, no grain sizes were given for the two materials. In a more recent study Yu et al. [55] presented polarization curves for polycrystalline (grain size 1 mm) and nanocrystalline (grain sizes 250 and 45 nm) copper in 0.1 M NaCl solution (Fig. 37.8). Again, the corrosion behavior of copper was not significantly changed by grain size reduction.

F. INDUSTRIAL APPLICATIONS AND OUTLOOK The results of the corrosion studies over the past two decades on nanocrystalline nickel, cobalt, zinc, copper, and several of

526


2.0

Potential (VSCE)

1.5

PC ~1000nm UFG ~250nm NC ~45nm

1.0 0.5 0.0 -0.5 -1.0 -1.5 10-10 10-9 10-8 10-7 10-6 10-5 10-4 10-3 10-2 10-1 Current Density (A/cm2)

FIGURE 37.8. Potentiodynamic polarization curves of electrodeposited copper foils with varying grain sizes in 0.1 M NaCl (pH 6.5).

their alloys made by electrodeposition can be summarized as follows: (i)

(ii)

(iii)

(iv)

(v)

Contrary to earlier concerns, the high-grain-boundary and triple-junction densities in nanocrystalline metals do not compromise their corrosion properties. The general shapes of polarization curves obtained in various solutions are not strongly affected by grain size reduction down to the nanometer range. Materials that show passivity in the polycrystalline form usually also exhibit passivity for the nanocrystalline structure. However, with decreasing grain size some changes in specific electrochemical parameters are observed, such as corrosion potential (Ecorr), corrosion current density (icorr), passivation potential (Ep), maximum current density (im), and passive current density (ip). For systems that show clear passivity, different studies found enhanced, reduced, or very similar passive current densities (in comparison with polycrystalline materials), depending on the system. A few discrepancies do exist in different investications for some materials (e.g., for the case of Cu) which would require further studies. Most likely there were some differences in the nanocrystalline and polycrystalline materials used in different studies, such as impurities or crystallographic texture. The structure and composition of the passive layer does depend on the grain size of the material. This can be understood on the basis of the high defect concentration (i.e., grain boundaries and triple junctions) intersecting the free surface of nanomaterials. Impurities also have a strong effect on the nature of the passive layer. Reduced grain size was observed in some studies to enhance the hydrogen evolution reaction.

(vi) Metals which are prone to intergranular attack (e.g, nickel, cobalt) can benefit enormously from grain size reduction because corrosion attack in nanomaterials is more evenly spread out over the entire surface instead of being concentrated in narrow channels along grain boundaries and triple junctions. The good corrosion performance of nanocrystalline nickel observed in the early studies [28, 29] was crucial in the development of the first large-scale industrial application of structural nanomaterials in the world: the so-called electrosleeve technology [35, 36]. This technology is essentially an in situ repair technology for nuclear steam generator tubing initially developed for Canadian CANDU reactors and later modified for other reactor types. In this repair approach nuclear steam generator tubing which was prematurely compromised by intergranular attack, intergranular stress corrosion cracking, and other forms of corrosion was repaired by the application of an electroformed nanocrystalline Ni–0.3% P microalloy sleeve on the inside of the tubes with a thickness of about 1 mm. The purpose of the sleeve was to (i) seal through-wall cracks to prevent leakage of radioactive water from the core of the reactor and (ii) restore the structural integrity of the tubes. While nanocrystalline nickel microalloy provided the necessary strength, it was the excellent corrosion resistance of nanocrystalline nickel which ultimately made this technology possible. The main purpose of using about 0.3% P in this microalloy as an alloying element was to achieve the long-term thermal stability of the sleeve material in a reactor environment operating at 280 C and a pH of 11.5. The phosphorus retards grain boundary mobility through a solute drag effect. The electrosleeve process [56, 57] has been successfully implemented in a CANDU unit in 1994 and a pressurized water reactor (PWR) in 1999. It has been incorporated by the American Society of Mechanical Engineers (ASME) as a standard procedure for pressure tubing repair [58]. Following this early application of electrodeposited nanocrystalline materials numerous applications have been developed over the past 10–15 years. Other applications are currently still at the research and development stage. Table 37.6 summarizes some of these applications, many of which require good to excellent corrosion resistance. For further details on the various applications the reader is referred to references [12–16] and [21–23]. Two of the most successful metallic coating systems used on many manufactured products are chromium, produced by electrodeposition from hexavalent chromium solutions, and cadmium, usually plated from alkaline, cyanide, acid sulfate, or acid fluorobate solutions. Chromium is extensively used for providing surfaces with wear, abrasion, and corrosion resistance or for decorative purposes, while cadmium is mainly used for corrosion resistance on many mass-finished products. However, considerable health risks are associated

REFERENCES

TABLE 37.6.

Applications of Electrodeposited Nanocrystals

Application

Materials

Armor laminates Battery grids Catalysts for H2 evolution Corrosion resistant coatings Chromium-replacement coatings Electromagnetic shielding Electronic connectors Electrosleeve Foil for printed circuit boards Free-standing soft magnets Hard-facing applications Magnetic recording heads Microelectromechanical systems Self-lubricating coatings Shaped charge liners Structural applications Transformer core materials Wear-resistant coatings

Ni, Fe, Co, Ni–Fe Pb Ni–Mo Ni, Ni–P, Zn–Ni, Co, Co–P Co–P Ni–Fe Ni Ni–P Cu Ni, Co, Ni–Fe, Co–Fe Ni, Ni–SiC, Ni–Al2O3, Co, Co–P Ni–Fe Ni, Co, Ni–Fe Ni–MoS2, Ni–BN Cu Ni, Co, Ni–Fe Ni–Fe, Co–Fe Ni, Ni–SiC, Ni–P, Co, Co–P

with both coating systems. Hexavalent chromium baths have been shown to enhance the risk of lung and nose cancer while exposure to cadmium can result in acute respiratory and gastrointestinal effects as well as chronic lung and kidney disease. For this reason several more benign coating systems have been developed to replace chromium and cadmium [15]. Various nanocrystalline-based alloys (e.g., Co–P, Co– Fe–P) are available for chromium replacement, while the best choices for cadmium replacement are Zn–Ni-type nanodeposits. The outlook for applications of nanocrystalline metals in applications requiring corrosion resistance is excellent. It is the unique combination of the overall good corrosion performance with outstanding mechanical properties which make this relatively new class of materials attractive to many industries.

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6. M. Naka, K. Hashimoto, and T. Masumoto, Corrosion, 36, 679 (1980). 7. J. C. Turn and R. M. Latanision, Corrosion, 39, 271 (1983). 8. S. J. Thorpe, B. Ramaswami, and K. T. Aust, J. Electrochem. Soc., 135, 2162 (1988). 9. P. Bragagnolo, Y. Waseda, G. Palumbo, and K.T. Aust, MRS. Int. Mtg. Adv. Mat., 4, 469 (1989). 10. U. Erb and A.M. El-Sherik, U.S. Patent No. 5,352,266, 1994. 11. U. Erb, A.M. El-Sherik, C. K. S. Cheung, and M. J. Aus,U.S. Patent No. 5,433,797, 1995. 12. C. K. S. Cheung, D. Wood, and U. Erb, in Processing and Properties of Nanocrystalline Materials, C. Suryanarayana et al., (Eds.), The Minerals, Metals and Materials Society, Warrendale, PA, 1996, p. 479. 13. G. Palumbo, F. Gonzalez, K. Tomantschger, U. Erb, and K.T. Aust Plat., Surf. Fin., 90(2), 36 (2003). 14. U. Erb, K. T. Aust, and G. Palumbo, in Nanostructured Materials, 2nd ed., C. C. Koch, (Ed.), William Andrew Publ., Norwich, NY, 2007, 235. 15. G. Palumbo, J. L. McCrea, and U. Erb, in Encyclopedia of Nanoscience and Nanotechnology, Vol. 1, H. S. Nalwa (Ed.), American Scientific Publ., Stevenson Ranch, CA, 2004, p. 89. 16. U. Erb, in CRC Materials Processing Handbook, J. R. Groza, E. J. Lavernia, J. F. Shackelford, and M. T. Powers,(Eds.), CRC Press, Boca Raton, FL, 2007, p. 22–1. 17. A. M. El-Sherik and U. Erb, J. Mater. Sci., 30, 5743 (1995). 18. R. T. C. Choo, A. M. El-Sherik, J. Toguri, and U. Erb, J. Appl. Electrochem., 25, 384 (1995). 19. H. Natter and R. Hempelmann, Z. Phys. Chem., 222, 319 (2008). 20. L. P. Bicelli, B. Bozzini, C. Mele, and L. D’Urzo, Int. J. Electrochem. Sci., 3, 356 (2008). 21. Erb U, Palumbo G, Aust KT, in Nanostructured Films and Coatings, G.M. Chow et al. (Eds.), NATO Science Series, 3 High Technol., 78, 11 (2000). 22. U. Erb, G. Palumbo, D. H. Jeong, S. H. Kim, and K. T. Aust, in Processing and Properties of Structural Nanomaterials, L. L. Shaw et al., (Eds.), The Minerals, Metals and Materials Society, Warrendale, PA, 2003, p. 109. 23. U. Erb, K. T. Aust, G. Palumbo, J. L. McCrea, and F. Gonzalez, in Processing and Fabrication of Advanced Materials IX, T. S. Srivatsan, et al., (Eds.), ASM International, Materials Park, OH, 2001, p. 253. 24. G. Palumbo, S. J. Thorpe, and K. T. Aust, Scripta Metall., 24, 1347 (1990). 25. A. M. El-Sherik, U. Erb, G. Palumbo, and K. T. Aust, Scripta Metall., Mater., 27, 1185 (1992). 26. N. Wang, Z. Wang, K. T. Aust, and U. Erb, Mater. Sci. Eng., A237, 150 (1997). 27. D. H. Jeong, K. T. Aust, U. Erb, G. Palumbo, Scripta Mater., 44, 493 (2001). 28. R. Rofagha, R. Langer, A. M. El-Sherik, U. Erb, G. Palumbo, K. T. Aust, Scripta Metall. Mater. 25, 2867 (1991). 29. R. Rofagha, R. Langer, A. M. El-Sherik, U. Erb, G. Palumbo, and K. T. Aust, Mater. Res. Soc. Symp. Proc., 238, 751 (1992).

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30. R. Rofagha, S. J. Splinter, U. Erb, N. S. McIntyre, Nanostr. Mater, 4, 69 (1994). 31. S. Wang, R. Rofagha, P. R. Roberge, and U. Erb, Electrochem. Soc. Proc., 95-8, 244 (1995). 32. P. T. Tang, T. Watanabe, J. E. T. Andersen, and G. BechNielsen, J. Appl. Electrochem., 25, 347 (1995). 33. A. M. El-Sherik and U. Erb, Plat. Surf. Fin., 82(9), 85 (1995). 34. S. H. Kim, K. T. Aust, U. Erb, and F. Gonzalez, in Proc. American Electroplaters and Surface Finishers Society (AESF) SUR/FIN 2002, AESF, Orlando, FL, p. 225 35. F. Gonzalez, A. M. Brennenstuhl, G. Palumbo, U. Erb, and P. C. Lichtenberger, Mater. Sci. For., 225–227, 831 (1996). 36. G. Palumbo, F. Gonzalez, A. M. Brennenstuhl, U. Erb, W. Shmayda, and P. C. Lichtenberger, Nanostr. Mater, 9, 737 (1997). 37. M. R. Zamanzad-Ghavidel, K. Raeissi, and A. Saatchi, Mater. Lett., 63, 1807 (2009). 38. C. Gu, J. Lian, J. He, Z. Jiang, and Q. Jiang, Surf. Coat. Technol., 200, 5413 (2006). 39. R. Rofagha, U. Erb, D. Ostrander, G. Palumbo, and K. T. Aust, Nanostr. Mater., 2, 1 (1993). 40. S. J. Splinter, R. Rofagha, N.S. McIntyre, and U. Erb, Surf. Interf. Anal., 24, 181 (1996). 41. L. Benea, P. L. Bonora, A. Borello, and S. Martelli, Wear, 249, 995 (2002). 42. X. Peng, Y. Zhang, J. Zhao, and F. Wang, Electrochim. Acta, 51, 4922 (2006). 43. S. H. Kim, K.T. Aust, U. Erb, F. Gonzalez, and G. Palumbo, Scripta Mater., 48, 1379 (2003).

44. S. H. Kim, T. Franken, G. D. Hibbard, U. Erb, K. T. Aust, and G. Palumbo, J. Metast. Nanostr. Mat., 15–16, 643 (2003). 45. A. Aledresse and A. M. Alfantazi, J. Mater. Sci., 39, 1523 (2004). 46. H. Jung and A. M. Alfantazi, Electrochim. Acta, 51, 1806 (2006). 47. M. Saito, K. Yamada, K. Ohashi, Y. Yasue, Y. Sogawa, and T. Osaka, J. Electrochem. Soc., 146, 2845 (1999). 48. K. M. S. Youssef, C. C. Koch, and P. S. Fedkiw, Corros. Sci., 46, 51 (2004). 49. A. M. Alfantazi and U. Erb, Corrosion, 52, 880 (1996). 50. S. Tao and D. Y. Li, Nanotechnology, 17, 65 (2006). 51. B. Yu, P. Woo, and U. Erb, Scripta Mater., 56, 353 (2007). 52. J. C. Hamilton, J. C. Farmer, and R. J. Anderson, J. Electrochem. Soc., 133, 739 (1986). 53. S. T. Mayer and R. H. Muller, J. Electrochem. Soc., 139, 426 (1992). 54. J. K. Yu, E. H. Han, L. Lu, and X. J. Wei, Mater. Sci., 40, 1019 (2005). 55. B. Yu, P. Woo, and U. Erb, to be published. 56. G. Palumbo, P. C. Lichtenberger, F. Gonzalez, and A. M. Brennenstuhl, U.S. Patent Nos. 5,527, 445, 5,516,415, and 5,538,615, 1996. 57. I. Brooks, G. Palumbo, F. Gonzalez, A. Robertson, K. Tomantschger and K. Panagiotopoulos, in American Electroplaters and Surface Finishers Society (AESF) SUR/FIN 2003 Proc., 2003, Orlando, FL, p. 721. 58. ASME Code Case 96-189-BC96-206, Case N-569-Section XI, Division 1: Alternative Rules for Repair by Electrochemical Deposition of Class 1 and 2 Steam Generator Tubing, 1996.

38 CORROSION OF SHAPE MEMORY AND SUPERELASTIC ALLOYS L. E. EISELSTEIN Exponent-Failure Analysis Associates, Inc., Menlo Park, California

A. Introduction A1. What are shape memory and superelastic alloys? A2. Shape memory and superelastic alloy classes B. NiTi-based shape memory alloys B1. NiTi passive layer B2. NiTi uniform corrosion B3. NiTi crevice and pitting corrosion B3.1. Pitting of NiTi B3.2. Crevice corrosion of NiTi B4. Stress corrosion cracking and hydrogen effects of NiTi B4.1. Hydrogen effects B4.2. Stress corrosion cracking of NiTi B5. NiTi galvanic corrosion B6. NiTi corrosion fatigue B7. NiTi fretting corrosion B8. NiTi high-temperature oxidation B9. Corrosion of NiTi medical devices B9.1. In vivo humans B9.2. In vivo animals B9.3. In vivo cell culture B9.4. Quality assurance testing for long-term NiTi medical device implants B9.5. Nickel release B9.5. Nickel release B9.6. Effect of synthetic physiological solutions on pitting C. b-Ti: Ni-Free Ti-based shape memory alloys D. Cu-based shape memory alloys D1. Cu-based SMA pitting D2. Cu-based SMA dezincification D3. Cu-based SMA intergranular attack and stress corrosion cracking E. Fe-based shape memory alloys F. Magnetic shape memory and shape memory alloys other than titanium based F1. Magnetic shape memory References

A. INTRODUCTION Shape memory and superelastic alloys are alloys that exhibit thermally recoverable strain (shape memory effect), pseudoelastic (superelastic behavior), or both. This chapter discusses what is known about the corrosion behavior of such alloys. However, most of the information on the corrosion behavior of such materials is focused on the 50 at % nickel–50 at % titanium (atomic percent, at %) (50 at % Ni–50 at % Ti) alloy composition since this material has extensive use in biomedical applications [1–12]. The corrosion data on other shape memory or superelastic materials will be discussed when available. We first describe what makes these alloys unique and what they are typically used for followed by a general description of the various alloy classes. There are also ceramic and polymer materials that exhibit shape memory; however, they are not discussed in this chapter [13–15]. A1. What Are Shape Memory and Superelastic Alloys? Shape memory and superelastic alloys have unique mechanical properties that make them useful, or even indispensable, for some applications. These properties are known as the shape memory effect and the superelastic or pseudoelastic € effect. In 1931 Olander was the first to point out that the gold–cadmium (Au–Cd) alloy had rubberlike characteristics [16–19]. This “rubberlike behavior” is still not well understood, but as far as the mechanical response, it appears to be superelastic. But unlike normal superelastic alloys in which this effect occurs as a result of stress-induced martensite in an austenitic matrix, the superelastic behavior of these rubberlike materials occur in the martensitic state [18]. The shape


529

530

CORROSION OF SHAPE MEMORY AND SUPERELASTIC ALLOYS

memory effect (thermoelastic behavior of the martensite phase) was described more fully in an Au–Cd alloy in 1951 [20] and then in indium–thallium (In-Tl) in 1952 [21]. The discovery of similar behavior in equiatomic nickel– titanium (NiTi) in the early 1960s was the beginning of commercial applications for such materials [22, 23]. Materials that exhibit crystallographically reversible martensitic transformation (or thermoelastic martensitic transformation) are generally called shape memory alloys (SMAs) [24, 25]. Such transformations can be induced by the application of force or straining or from changes in temperature. The martensite transformation is accompanied by a large diffusionless shear-like deformation associated with structural change. This deformation generally amounts to about 20 times the elastic deformation. The martensite is deformable, and it can be induced from the parent austenite phase by loading. A large deformation induced in shape memory alloys can be recovered by heating to temperatures above the reverse transformation finish temperature (Af) after unloading (the shape memory effect) or simply by unloading at temperatures above Af (the pseudoelasticity or superelasticity effect) [25]. These transformations can occur in a wide variety of alloys such as Ag–Cd, Cu–Al–Ni, NiTi, Cu–Sn, and Cu–Zn [24]. Rubber-like superelastic behavior has also been observed in Au–Cu–Zn, Cu–Al–Ni, Cu–Zn–Al, and Au–Cd alloys [18]. A2. Shape Memory and superelastic Alloy Classes The three main commercial types of shape memory alloys are the copper–zinc–aluminum–nickel (Cu–Zn–Al–Ni), copper–aluminum–nickel (Cu–Al–Ni), and nickel–titanium (NiTi) alloys although there are many other SMA families. Examples of various SMAs are shown in Table 38.1. Although shape memory and superelastic alloy classes could be described any number ways, they are broken down into five general classes in this chapter. The five classes we will consider will be NiTi, the founding member class; Ni-free Ti-based alloy composition class; Cu-based, Fe-based, and magnetic shape memory, and shape memory alloys other than Ti-based alloys, a catchall class. Generally only the NiTi, Cu-based, and ferrous families are currently considered materials of commercial interest [25].

B. NiTi-BASED SHAPE MEMORY ALLOYS The NiTi shape memory alloy was first described by Buehler in the early 1960s [22, 23]. NiTi [50 at % Ni–50 at % Ti, also known as nitinol, since Buehler worked for the Naval Ordinance Laboratory (NiTiNOL)] is the most commonly used superelastic and SMA; thus, it is the alloy composition for which we have the most corrosion data. The Ti–Ni SMAs have become one of the most important metallic biomedical

TABLE 38.1. Various Shape Memory Alloy Compositions Alloy (Typical Composition, wt %)

References

NiTi Based (55–58%)NiTi Ni–(45–46%)Ti–(2 mm in diameter) are usually natural materials generated from rocks, soils, and so on. Because of their size, these can be filtered with high efficiency from the air in indoor environments. Coarse particles usually are not corrosive but can be abrasive or introduce a high resistance into a contact or connector. Fine particles (60% RH) and ionic contamination, as is common in many urban, suburban, and industrial environments [4, 5]. Printed circuit boards made from glass fiber–reinforced epoxy may be susceptible to failure caused by the formation of conductive anodic filaments [16, 17]. The first step in this failure mode is the formation of a poor glass/epoxy bond or the delamination of this interface. At sufficiently high relative humidity, moisture is adsorbed by the glass along this interface, reducing the insulation resistance of the circuit board. In the presence of an applied voltage, the positively biased copper conductor becomes the anode and is oxidized in an electrochemical cell. Water is also oxidized at the anode, forming hydronium ions, 2H2 O ! O2 þ 4H þ þ 4e

ð40:3Þ

At the negatively biased copper conductor, water is reduced to form hydroxyl ions, 2H2 O þ 2e ! H2 þ 2OH

ð40:4Þ

Consequently, a gradient in pH develops between the positively and negatively biased conductors. Copper ions formed at the anode migrate along this concentration gradient until they reach a neutral pH at which they precipitate, forming a filament along the epoxy–glass interface. Eventually, a conductive bridge is formed between the two copper conductors. Other degradation mechanisms include tarnishing by reactive gases, for example, moist air and sulfur-containing gases, corrosion by sulfur in packaging materials, rubber, and plastics, corrosion by reactive compounds emitted by organic materials, such as adhesives and coatings, during curing or aging [18], galvanic corrosion between dissimilar metals, and electrolytic corrosion caused by the presence of bromide

OTHER FAILURES AND FAILURE MECHANISMS

from the decomposition of fire retardants in plastic boards (see Section C).

E. HYBRID INTEGRATED CIRCUITS AND MULTICHIP MODULES On HICs and multichip modules, layered metallizations may be used instead of aluminum to interconnect individual ICs and other components. Common layered metallizations include chromium/copper, titanium/copper, titanium/palladium/gold, and titanium/palladium/copper/nickel/gold. Silver, tin, and lead are also found in some metallizations. Minimum line widths and line spacings may be 5 mm or less. The corrosion problems associated with HICs and multichip modules are similar to those described previously for integrated circuit and printed circuit boards. The multilayered metallizations make these circuits more prone to galvanic corrosion, particularly during manufacture.

F. CONTACTS AND CONNECTORS Contacts and connectors may be characterized by their contact force and the voltage across them. If the contact force is sufficiently high, most tarnish films are penetrated upon closing and good contact is made. Similarly, a sufficiently high voltage across the contact surfaces will destroy most tarnish films and produce good contact. However, for low-force, low-voltage contacts a tarnish film of a-FeOOH > g-FeOOH; b-FeOOH, amorphous rust g-FeOOH, a-FeOOH, Fe3O4; amorphous

Reproduced with permission from [18].

Long-term atmospheric exposure tests have been conducted at a large number of locations around the world. Because it is not possible to present in this chapter all the data that have been obtained, and because data obtained at one location are not helpful for corrosion engineering at other locations, only a list of major exposure programs is given in Table 43.3. A survey of the data obtained in these programs shows that the average values of corrosion penetration in the first 10 years are in the following ranges: Industrial atmospheres 0:1--0:5 mm Marine atmospheres 0:3--0:8 mm Rural atmospheres 0:05--0:2 mm The method of exposure affects the corrosion rate. In most of the tests, specimens in the form of plates or sheets were exposed to the atmosphere at an angle of 30 from the horizontal with the racks facing south. The corrosion rate was measured by mass loss caused by corrosion of both surfaces of the specimen and, using the specific density, converted to the decrease in thickness from one surface so that the value was the average of corrosion of the skyward and groundward surfaces. In polluted atmospheres, corrosion that occurs on the groundward surfaces is greater than on the skyward surfaces, because the former surfaces do not benefit from the washing action of rain and the quick drying action of the sunshine that enhance the protective quality of the corrosion product films. It has been shown [60] that for steel specimens exposed in an industrial atmosphere on the 30 rack the corrosion that occurred on the skyward and groundward surfaces was 40 and 60%, respectively, of the total corrosion. In some programs, specimens were exposed vertically. LaQue [61] pointed out that the vertical exposure has the disadvantage of poorer reproducibility because a slight variation in the position of the specimen from the vertical greatly affects the washing and drying effects. Galvanic corrosion data based on atmospheric exposure tests were surveyed by Kucera and Mattsson [62]. In rural and industrial atmospheres, no risk of galvanic corrosion of

584

CARBON STEEL—ATMOSPHERIC CORROSION

TABLE 43.3. Major Atmospheric Exposure Tests in Various Countries Maximum Test Period (years)

References

Country

Organization

Locationa

Started In

International

ISO

12 countries (47 sites)

1987

8

–b

Australia

Commonwealth Scientific and Electricity Trust of South Australia (ETSA)/CSIRO

Melbourne (299 sites) South Australia (66 sites) South Australia (475 sites)

1979 1990 1991

2 2 2

23 24 24

Brazil

Campanhia Siderurgica Natuinal (CSN)

Volta Redonda (U/I) Volta Redonda (R) Restinga da Marambaia (M)

1972

16

25

Centro de Pesquisas e Deservolvimento Leopoldo A. Miguez de Mello (CENPES) of Petrobras

Manaus (E)c Belem (E)c Fortaleza (M) Aracaju (5 sites, M) Ajacaju (R) Madre de Deus (M) Betim (I) Rio de Janeiro (M) Cubat~ao (I) S. Mateus (I) Canoas (l) Brasilia (U)

1975

2

25

Canada

CANMET

Arctic & Southern Canada (77 sites)

1978

1

26

Czechoslovakia

Akimov State Research Inst. of Material Protection

Prague (U) Usti nad Labem (I) Hurbanovo (R/I) Kopisty (I)

Late 1960s

5

27

Finland

Technical Research Center (VTT)

Otaniemi (R) Helsinki (U/I) Harmaja (M) Otaniemi (R) Koski (R) Haravalta (I)

1968 1975

7 4

28 28

Germany

Staatlichen Materialpr€ufungsamt, Berlin Iron and Steel Inst.

Berlin (R) Dortmunt (I) H€ ornum (M) Olpe (R) M€ ulheim (I) Cuxhaven (M)

1914

4

29

1956

8

30

Vereins Deutscher Eisenh€uttenleute

Duisburg (I) Gelsenkirchen (I) M€ ulheim (I) Cuxhaven (M) Olpe (R)

1962

16

31, 32

(about 1970)d

8

32

Research Group on Corrosion Protection of Steel Structures

Obihiro (R) Wajima (M) Tokyo (I) Kawasaki (I) Omaezaki (M) Takayama (R) Makurazaki (M)

1960

5

5

Steel Manufacturers

38 sites in Japan

1960s

12 (max)

33

Public Works Research Inst.

41 sites in Japan (sheltered)

1981

9

14

Japan

585

CORROSION DATA

TABLE 43.3.

(Continued )

Country

Organization

Locationa

Started In d

Maximum Test Period (years)

References

4

34

Latin America

Project on “Ibero-American Map of Atmospheric Corrosiveness”

72 sites

(about 1990)

New Zealand

Building Research Assoc. of New Zealand

168 sites in NZ

1987

1

35

Norway

VERITAS and Norwegian Inst. for Air Research

Sogn (R) Voss (R) Bergen-Bergens Tidende (U) Bergen-VERITAS (U) Bergen-Stend (R) Bergen-Fredriksberg (U) Bergen-Minde (U) Stord 50 (M/R)

1970

8

36

Panama

Naval Research Laboratory

Cristobal (M) Miraflores (Semi-U)

by 1958

16

37, 38

Russia (Formerly USSR)

Inst. of Physical Chemistry

Moscow (I) Zvenigorod (R) Batumi (M/U)

1968

10

39

Scandinavia

Scandinavian Council for Applied Research (NORDFORSC)

32 sites (Norway 19, Sweden 6, Finland 5, Denmark 2)

1975

8

40

Singapore

National University of Singapore

Singapore (M) A raft (M)

1985/86

12 weeks

41

Spain

Ciudad University

South (16 sites) Northwest (21 sites) Center (34 sites)

1976/83

1

42

University of Barcelona

Catalonia (42 sites, R, U, I, M)

(1980s)d

Centro Nacional de Investigaciones Metalurgicas

El Escorial (R) Madrid (U) Zaragoza (U) Bilbao (I) Barcelona (M) Cadiz (M) Cabo Negro (M) Alicante (M)

1976

Sweden

Swedish Corrosion Inst.

Ryda (R) Stockholm (U) Borregaard (I) Bohus-Malm€ on (M)

Taiwan

China Steel Corp.

United Kingdom

3

43

13–16

44

(1980s)d

5

45

Hsinchu (R/U) Kaohsuiung (M) Hsiao Kang (I)

(1980s)d

8

46

Iron Steel Inst.

Calshot (M) Dove Holes Tunnel Llanwrtyd Wells (R) Motherwell (I) Redear (M/I) Sheffield (I) Woolwich (I)

1928

5

47

BISRA

Sheffield (I)

1937

5

48 (continued )

586


TABLE 43.3. (Continued ) Maximum Test Period (years)

Country

Organization

Locationa

Started In

United States

ASTM (Committee A-5)

Pittsburgh, PA (I) Fort Sheridan, IL (R) Annapolis, MD (M)

1916

7.3 12 >30

49

U.S. Steel

Kearney, NJ (I) Kure Beach, NC (M, 250 m)

1938 1940

11.5 7.5

50

INCO

Kure Beach, NC (M, 250 m) Block Island, RI (M) Bayonne, NJ (I)

1941

15.5 17.1 18.1

51

U.S. Steel

Kearney, NJ (I) South Bend, PA (R) Kure Beach, NC (M)

1942

15.5

52

ASTM (Committee A-7) (malleable)

Kure Beach, NC (M, 25 m) Newark NJ (I) Point Reyes, CA (M) State College, PA (R) E. Chicago, IN (I)

1958

12

53, 54

Pacific Gas & Electric Co. Steel Founders’ Society of America

Northern CA (15 sites)

1956

5

55

Kure Beach, NC (M, 25, 250 m) East Chicago, IN (I)

1953

7

56, 57

Bethlehem Steel

Kure Beach, NC (M, 250 m) Saylorsburg (R) Bethlehem, PA (I)

1968

16

58

USX

Kearney, NJ (U/I) Kure Beach, NC (M, 250 m) Potter County, PA (R)

Early 1970s

16

59

References

12

a

Rural R, Urban U, industrial I, Marine M. To be published. c Equatorial. d Not stated in reference and estimated. b

carbon steel exists when the steel is in contact with Pb, Zn, Al, Mg, or weathering steel. Increased corrosion results if carbon steel is coupled to stainless steels, Cu, Ni, anodized Al, Sn, or Cr. In marine atmospheres, the effect is essentially the same, except that contact with anodized Al and Cr is not harmful. Unlike the situation in aqueous environments, galvanic corrosion in the atmosphere is usually restricted to a narrow zone in the anodic metal, and the galvanic effects do not extend over several millimeters from the line of contact [63].

REFERENCES 1. International Organization for Standardization (ISO), Corrosion of Metals and Alloys—Corrosivity of Atmospheres— Classification, ISO 9223-1992, ISO, Geneva, Switzerland, 1992.

2. G. Schikorr, Werkst. Korros., 15, 457 (1964). 3. H. Schwarz, Werkst. Korros., 16, 93, 208 (1965). 4. American Society for Testing and Materials (ASTM) Task Force on the Calibration of Atmospheric Corrosivity, ASTM STP 435, ASTM, Philadelphia, PA, 1968, p. 360. 5. K. Horikawa, S. Takiguchi, Y. Ishizu, and M. Kanazashi, Boshoku Gijutu (Corros. Eng.), 16, 153 (1967). 6. International Organization for Standardization (ISO), Corrosion of Metals and Alloys—Corrosivity of Atmospheres— Measurement of Pollution, ISO 9225-1992, ISO, Geneva, Switzerland, 1992. 7. Rikujo Tekkotsu Kozobutsu Boshoku Kenkyu Kai, Memorial Lecture, 1967 Award for Corrosion Protection, Japan Society for Promotion of Science, Tokyo, Japan. 8. Rikujo Tekkotsu Kozobutsu Boshoku Kenkyu Kai, Boshoku Gijutu (Corros. Eng.), 22, 106 (1973).

REFERENCES

9. Study on Metallic Corrosion by Atmospheric Pollutants, 2nd Report, Tokyo Metropolitan Research Institute for Environmental Protection, Tokyo, Japan, 1978. 10. International Organization for Standardization (ISO), Corrosion of Metals and Alloys—Corrosivity of Atmospheres— Guiding Values for the Corrosivity Categories, ISO 92241992, ISO, Geneva, Switzerland, 1992. 11. International Organization for Standardization (ISO), Corrosion of Metals and Alloys—Corrosivity of Atmospheres— Determination of Corrosion Rate of Standard Specimens for the Evaluation of Corrosivity, ISO 9226-1992, ISO, Geneva, Switzerland, 1992. 12. Provided by Y. Togawa, based on Y. Togawa, Bousei Kanri (Rust Prevention and Control), 37(2), 7 (1993). 13. R. Passano,Washington Regional Meeting ASTM, March 7, 1934, ASTM, Philadelphia, PA, cited in H. H. Uhlig and R. W. Revie, Corrosion and Corrosion Control, 3rd ed., Wiley, New York, 1985, p. 167. 14. Public Works Research Institute, Kozai Club and Japan Association of Steel Bridge Construction, Joint Research on Application of Weathering Steel to Bridges, Report XVIII, Tokyo, Japan, 1993. 15. I. Matsushima and T. Ueno, Corros. Sci., 11, 129 (1971). 16. Y. Shimizu, K. Tanaka, and T. Nishimura, Zairyo-to-Kankyo (Corros. Eng.) 44, 436 (1995). 17. P. Keller, Werkst. Korros., 18, 865 (1967). 18. T. Misawa, Boshoku Gijutu (Corros. Eng.), 32, 657 (1983). 19. U. R. Evans, Trans. Inst. Metal Finishing, 37, 1 (1960). 20. U. R. Evans and C. A. J. Taylor, Corros. Sri., 12, 227 (1972). 21. M. Stratmann, K. Bohnenkamp, and H-J. Engell, Cor. Sci., 23, 969 (1983). 22. J. D€unnwald and A. Otto, Corros. Sci., 29, 1167 (1989). 23. J. F. Moresby, F. M. Reeves, and D. J. Spedding, in Atmospheric Corrosion, W. H. Ailor (Ed.), Wiley, New York, 1982, p. 745. 24. G. A. King, J. Kapetas, and D. Bates-Brownsword, Paper No. 106, CORROSION/94, NACE, 1994. 25. A. C. Dutra and R. de O. Vianna, in Atmospheric Corrosion, W. H. Ailor (Ed.), Wiley, New York, 1982, p. 755. 26. G. J. Biefer, Mater. Perform., 20(1), 16 (1981). akova, J. Vickova, and J. Honzak, in 27. D. Knotkova-Cer Atmospheric Corrosion of Metals, ASTM STP 767, S. W. Dean, Jr., and E. C. Rhea (Eds.), ASTM, Philadelphia, PA, 1982, p. 7. 28. T. Hakkarainen and S. Yl€asaari, Atmospheric Corrosion, W. H. Ailor (Ed.), Wiley, New York, 1982, p. 787. 29. O. Bauer, Stahl Eisen, 41, 37, 76 (1921). 30. E. Brauns and U. Kalla, Stahl Eisen, 85, 406 (1965). 31. W. Schwenk and H. Ternes, Stahl Eisen, 88, 318 (1968). 32. G. Burgmann and D. Grimme, Stahl Eisen, 100, 641 (1980). 33. Committee on Steel Corrosion Survey, Kozai Club, Committee Report, A Survey on Corrosion and Corrosion Protection of Steel, Kozai Club, Tokyo, Japan, 1973, p. 18; reproduced in Corrosion and Corrosion Protection Databook, I. Matsushima

34.

35. 36. 37. 38. 39. 40.

41. 42. 43.

44.

45.

46. 47. 48. 49. 50. 51. 52.

53.

54. 55.

56.

57.

587

(Ed.), Japan Society of Corrosion Engineering, Tokyo, Japan, 1995, p. 6. M. Morcillo, in Atmospheric Corrosion, ASTM STP 1239, W. W. Kirk and H. H. Lawson, (Eds.), American Society for Testing and Materials, Philadelphia, PA, 1995, p. 257. R. J. Cordner, Br. Corros. J., 25, 115 (1990). L. Atteraas and S. Haagenrud, in Atmospheric Corrosion, W. H. Ailor (Ed.), Wiley, New York, 1982, p. 837. C. R. Southwell, J. D. Bultman, and A. L. Alexander, Mater. Perform., 15(7), 9 (1976). C. R. Southwell and J. D. Bultman, in Atmospheric Corrosion, W. H. Ailor (Ed.), Wiley, New York, 1982, p. 943. Y. N. Mikhailovski and P. V. Strekalov, in Atmospheric Corrosion, W. H. Ailor (Ed.), Wiley, New York, 1982, p. 923. V. Kucera, S. Haagenrud, L. Atteraas, and J. Gullman, in Degradation of Metals in the Atmosphere, ASTM STP 965, S. W. Dean and T. S. Lee (Eds.), American Society for Testing and Materials, Philadelphia, PA, 1988, p. 264. S. K. Roy and K. H. Ho, Br. Corros. J., 29, 287 (1994). S. Feliu and M. Morcillo, Br. Corros. J., 22, 99 (1987). E. Brillas, J. M. Costa, and M. Villarrasa, Proceedings of the 11th International Corrosion Congress, Associazione Italiana di Metallurgia (AIM), Milan, Italy, 1990, vol. 2, p. 2. 79. M. Morcillo, J. Simancas, and S. Feliu, in Atmospheric Corrosion, ASTM STP 1239, W. W. Kirk and H. H. Lawson (Eds.), American Society for Testing and Materials, Philadelphia, PA, 1995, p. 195. E. Johansson and J. Gullman, in Atmospheric Corrosion, ASTM STP 1239, W. W. Kirk and H. H. Lawson (Eds.), American Society for Testing and Materials, Philadelphia, PA, 1995, Vol. 2, p. 2.79. Feng-I. Wei, Br. Corros. J., 26, 209 (1991). J. C. Hudson, J. Iron Steel Inst., 148, 161 (1943). J. C. Hudson and J. F. Stanners, J. Iron Steel Inst., 180, 271 (1955). ASTM Committee A-5, Proc. ASTM, 53, 110 (1953). C. P. Larrabee, Corrosion, 9, 259 (1953). H. R. Copson, Proc. ASTM, 60, 650 (1960). C. P. Larrabee and S. K. Coburn, in Metallic Corrosion (Proceedings of the 1st International Congr. Metallic Corrosion), Butterworths. London, 1971, p. 276. G. B. Mannweiler, in Metal Corrosion in the Atmosphere, ASTM STP 435, American Society for Testing and Materials, Philadelphia, PA, 1968, p. 211. C. McCaul and S. Goldspiel, in Atmospheric Corrosion, W. H. Ailor (Ed.), Wiley, New York, 1982, p. 432. H. E. Thomas and H. N. Alderson, in Metal Corrosion in the Atmosphere, ASTM STP 435, American Society for Testing and Materials, Philadelphia, PA, 1968, p. 83. C. W. Wieser, in Metal Corrosion in the Atmosphere, ASTM STP 435, American Society for Testing and Materials, Philadelphia, PA, 1968, p. 271. P. F. Wieser, in Atmospheric Corrosion, W. H. Ailor (Ed.), Wiley, New York, 1982, p. 453.

588


58. C. R. Shastry, J. J. Friel, and H. E. Townsend, in Degradation of Metals in the Atmosphere, ASTM STP 965, S. W. Dean and T. S. Lee (Eds.), American Society for Testing and Materials, Philadelphia, PA, 1988, p. 5. 59. S. K. Coburn, M. E. Komp, and S. C. Lore, in Atmospheric Corrosion, ASTM STP 1239, W. W. Kirk and H. H. Lawson (Eds.), American Society for Testing and Materials, Philadelphia, PA, 1995, p. 101.

60. C. P. Larrabee, Trans. Electrochem. Soc., 85, 297 (1944); cited in refs. [51] and [61]. 61. F. L. LaQue, Proc. ASTM, 51, 1 (1951). 62. K. Kucera and E. Mattsson, in Atmospheric Corrosion, W. H. Ailor (Ed.), Wiley, New York, 1982, p. 561. 63. I. L. Rozenfeld, Corrosion and Corrosion Protection of Metals, Izd. Metallurgija, Moscow, Russia, 1970, p. 120.

44 CARBON STEEL—CORROSION IN FRESHWATERS I. MATSUSHIMA* Maebashi Institute of Technology, Maebashi, Japan

A. Introduction B. Environmental factors B1. Natural freshwaters B2. Effect of water quality B2.1. General B2.2. Saturation index B2.3. pH B2.4. Chloride and sulfate ions B2.5. Bicarbonate B3. Effect of velocity C. Localized corrosion C1. General C2. Pitting under tubercles C3. Groove corrosion of electric resistance welded pipe C4. Localized corrosion at discontinuities in corrosion product films References

A. INTRODUCTION Corrosion of steel in freshwater proceeds electrochemically by the action of dissolved oxygen: Anodic reaction: Fe ! Fe2 þ þ 2e

Cathodic reaction:

Overall:

1 2 O2

þ H2 O þ 2e ! 2OH

Fe þ 12 O2 þ H2 O ! FeðOHÞ2

ð44:1Þ ð44:2Þ ð44:3Þ

Most natural freshwaters are air saturated, and the concentration of dissolved oxygen at ordinary temperatures is 8–10 ppm. For corrosion of steel to occur, dissolved oxygen must be supplied to the steel surface by diffusion through water and possibly through surface films on the steel surface that act as diffusion barrier layers. Because of its concentration and diffusion coefficient, the supply of dissolved oxygen is slow unless the velocity of water relative to the steel surface is sufficiently high. The corrosion reaction proceeds as rapidly as oxygen reaches the steel surface. Thus, the corrosion is under diffusion control of dissolved oxygen. In the absence of diffusion barrier layers on the surface, the theoretical corrosion current density, i (A/cm2), of steel in stagnant air-saturated freshwater can be calculated as follows: DnF C 10 3 ð44:4Þ i¼ d where D is the diffusion coefficient for dissolved oxygen in water (cm2/s), n is the number of electrons involved in the reaction, F is the Faraday constant (C/mol), d is the thickness of the diffusion layer (cm), and C is the concentration of dissolved oxygen (mol/L). Using D ¼ 2 10 5 cm2 =s; n ¼ 4; F ¼ 96; 500 C=mol; d ¼ 0:05 cm, and C ¼ 8=32 10 3 mol=L at 25 C; i ¼ 38:6 10 6 A=cm2 , which is equivalent to a corrosion rate of 0.45 mm/year [18 mils/year (mpy)]. In hard waters that contain high concentrations of calcium and bicarbonate, the natural deposition of calcium carbonate (CaCO3) on the steel surface provides an effective diffusion barrier to oxygen diffusion, greatly decreasing corrosion. In soft waters, the corrosion rate is higher than in hard waters, but it is lower than the theoretical maximum value because

*Deceased


589

590

CARBON STEEL—CORROSION IN FRESHWATERS

the corrosion product film formed on the surface acts to some extent as a diffusion barrier. The average corrosion rate of steel in stagnant air-saturated soft waters at ordinary temperatures is roughly 0.1 mm/year. Increase in oxygen concentration, velocity, and temperature, within limits, accelerates corrosion in soft water by increasing the supply of dissolved oxygen to the surface. When oxygen concentration is very high, for example, 12 mL/L or 17 ppm in distilled water at 25 C [1] or 25–35 ppm [2] in natural waters (e.g., when freshwater is saturated with a gas of high oxygen partial pressure), or water velocity is above a critical value, the corrosion rate drops to a low value. This decrease in corrosion rate is caused by passivation of the steel as a result of oxygen being supplied to the surface in excess of the amount that can be used for the cathodic reaction. More quantitative data regarding the effects of environmental factors are given in Section B. While the supply of dissolved oxygen controls the overall corrosion rate, the cathodic and anodic reactions do not necessarily occur uniformly over the surface. There are two types of corrosion cells. In macrogalvanic cells, the anodic and cathodic areas are macroscopic, and their locations are fixed, whereas in macrogalvanic cells, the anodic and cathodic sites are microscopic and move randomly with time. When macrogalvanic cells, such as differential aeration cells, are formed on the steel surface and the anodic reaction dissolution of steel takes place predominantly at some limited areas in amount equivalent to the cathodic reaction proceeding around the macroanode, the penetration rate at the anode can be high if the cathode–anode area ratio is large. The extent of the effective cathode area depends on the conductivity and geometry of the system. In waters of high conductance, macrogalvanic cells can operate over long distances. High corrosion rates of steel are associated with the formation of macrogalvanic cells of some kind, which may not have been anticipated or considered at the design stage for corrosion protection.

B. ENVIRONMENTAL FACTORS B1. Natural Freshwaters While corrosion of carbon steel proceeds by the action of dissolved oxygen, the rate of the reaction is affected by other species in water. Among the many species dissolved in natural freshwaters, calcium (Ca2 þ ), bicarbonate (HCO 3 ), and chloride (Cl ) are very important with respect to corrosion. The source of natural freshwaters is meteoric water that falls on the earth as rain, snow, sleet, or hail. Rainwater comes in contact with the atmosphere and is saturated with dissolved air. Depending on the kind of atmosphere to which the

rainwater is exposed, it also dissolves impurity gases, such as SOx, NOx, NH3, and HCl, and other matter from suspended atmospheric impurities such as sea salt and dust. When the precipitation comes into contact with the ground, some of it evaporates, but the rest flows or collects on the surface of the earth as surface water or sinks into the ground as groundwater. The water dissolves materials with which it comes in contact and therefore contains a variety of species, including Ca2þ , Mg2þ , soluble silica [e.g., þ þ 2 . (H2SiO3)n], HCO 3 , Na , K , Cl , and SO4 Carbon dioxide in water reacts with minerals containing CaCO3, such as limestone, marble, chalk, calcite, and dolomite, as follows: CaCO3 þ CO2 þ H2 O ! Ca2 þ þ 2HCO3

ð44:5Þ

Dissolved CO2 also reacts similarly with minerals containing 2þ and MgCO3 to form Mg2þ and HCO 3 . The contents of Ca 2þ Mg as equivalent parts per million CaCO3 are the calcium hardness and the magnesium hardness, respectively, and the sum is the total hardness. In natural waters, the concentration of bicarbonate (HCO3 ) corresponds to the methyl orange alkalinity (MO alkalinity; the equivalent per liter of titratable base to the methyl orange end point), commonly expressed as parts per million CaCO3. The values of Ca hardness and the MO alkalinity are the primary factors that determine the ease of protective CaCO3 film formation on a metal surface. The solubility of CaCO3 depends on the dissolved CO2 concentration. Air-saturated distilled water contains only 0.5 ppm CO2 at room temperature and can dissolve 53 ppm CaCO3. Rainwater, however, may contain up to 2 ppm CO2 because of the CO2 derived from organic matter in the atmosphere. The amount of CO2 contained in some natural waters, particularly groundwaters, is much higher due to the CO2 derived from the decay process of organic matter in the soil, and the Ca hardness and MO alkalinity may be over 200 ppm. At a given CO2 concentration, the amount of CaCO3 dissolution depends on the kind of geological strata to which the water is exposed and the contact time before the water is utilized. If the dissolution of CaCO3 and MgCO3 is the major source of the total hardness and the MO alkalinity, these two values should be almost identical. In some natural waters, however, the total hardness may be significantly higher or lower than the MO alkalinity. This difference can occur because Ca2 þ and Mg2 þ can derive from water-soluble minerals, instead of carbonates, without the action of CO2, for example, Ca2 þ from gypsum, alabaster or selenite and Mg2 þ from epsomite, kainite, picromerite, or loweite. Additional HCO 3 may be formed by the reaction of dissolved CO2 with some minerals other than CaCO3 and MgCO3, for example, albite (NaAlSi3O8). If the MO alkalinity of the water equals or exceeds the total hardness, all of

ENVIRONMENTAL FACTORS

TABLE 44.1.

River

591

Typical Water Qualities of River Watersa,b

Location

pH

MO Alkalinity (as CaCO3)

Hardnessc (as CaCO3) Total

Cl

Ca

SO42

SiO2

Total Dissolved Solids

(Concentration in ppm) Bandas Kerteh Yodo Vanda Enim Changjiang Indus Chaopraya Thachin Changjiang Donau Thames Red Deer Saskatchewan

Seria (Brunei) Kerteh (Malaysia) Sakai (Japan) Helsinki (Finland) Tanjung (Indonesia) Wuhan (China) Jamshoro (Pakistan) Bangkok (Thailand) Banglen (Thailand) Shanghai (China) Bratislava (Slovakia) London (UK) Red Deer (Canada) Saskatchewan (Canada)

5.3 6.2 6.7 7.5 8.3 7.2 8.0 6.8 7.5 8.0 7.8 8.4 8.0 7.3

1 3 24 38 141 99 92 107 150 102 178 215 218 380

8 15 39 50 73 75 125 126 150 163 233 262 290 604

1 7 29 31

1 10 15 12 12 17 15 109 35 153 23 19 6 476

53 73 100 98 169 252 180 352

7 1 22 22 18

6 10 6 8 29 5

28 50 40 33 27 37 298

26 20 10 9 6 7 14

78 145

904 1650

a

Data were supplied from the Kurita Water Industries database [3]. Water quality is subject to seasonal fluctuation. c Arranged in increasing total hardness. b

the hardness is attributed to carbonate hardness. If, on the other hand, the MO alkalinity is less than the total hardness, the carbonate hardness equals the MO alkalinity and the balance is the noncarbonate hardness. The results of analyses of dissolved matter in typical soft and hard river waters are listed in Table 44.1 [3]. Waters with a total hardness of < 100 ppm are soft, those with 100–150 ppm are slightly hard, and those with > 150 ppm are hard. Some waters are very hard, the total hardness being over 300 ppm. The sources of Cl in natural freshwaters are sea salt in the atmosphere, pollution of rivers by sewage and industrial effluents, some geological strata to which the water is exposed, and road deicing salts. The average concentration of Cl in rainwater is commonly 1 ppm, but near the seacoast, it can be an order of magnitude higher due to prevailing winds from the sea. The Cl content of river waters, about 8 ppm on average, varies in a wide range from a few parts per million to several hundred parts per million except in some unusually high cases. B2. Effect of Water Quality B2.1. General. The analytical water quality parameters that may affect corrosivity of natural freshwaters are pH, dissolved oxygen, Ca hardness, MO alkalinity, total dissolved solids (TDS), Cl , and SO42 . The values of pH, Ca hardness, MO alkalinity, and TDS are factors that determine the saturation index, which is the criterion of whether or not the CaCO3 diffusion-barrier film is formed on a metal

surface, as will be described in the following section (see Section B2.5). The values of pH and of dissolved oxygen concentration are the fundamental factors that affect the corrosivity of soft waters. They usually do not change the corrosion rate of steel in natural freshwaters, because the pH value remains within a certain range and most waters are air saturated. The ions Cl and SO42 are always present in natural freshwaters. They affect the penetration rate of localized corrosion by increasing conductivity, and they also affect the critical concentration of oxygen and critical water velocity above which passivation of steel occurs. B2.2. Saturation Index. For given values of calcium hardness, MO alkalinity and total dissolved salt concentration, a value of pH (pHs) exists at which the water is in equilibrium with solid CaCO3. The deposition of CaCO3 is thermodynamically possible when the pH of the water is higher than pHs. At equilibrium, K2 ¼

½H þ ½CO3 2 ½HCO3

ð44:6Þ

Ks ¼ ½Ca2 þ ½CO3 2

ð44:7Þ

½Ca2 þ ½HCO3 K2 ½H þ

ð44:8Þ

¼

592


Since log [H þ ] here is pHs and ½HCO3 ’ MO alkalinity (alk) in natural freshwaters,

Ks log Ca2 þ þ logðalkÞ pHs ’ log K2

ð44:9Þ

For a more thorough derivation of pHs, see the textbook by Uhlig and Revie [4], and for a new equation based on a rigorous model of CaCO3 saturation, see the paper by Pisigan and Singley [5]. The value pHmeasured pHs is called the saturation index, or the Langelier index, after Langelier [6], who established this index. A positive value of the Langelier index indicates a tendency for the protective CaCO3 film to form and a negative value for it not to form. In the presence of other dissolved salts, the increase of ionic strength depresses the activities of other ions and increases the values of log(Ks/K2) in Eq. 44.9 if Ks and K2 are based on concentrations rather than activities. On the other hand, the values of log(Ks/K2) decrease with temperature, making the saturation index more positive. In practice, the calculation of pHs is carried out on the concentration basis with additions of correction factors for ionic strength and temperature to Eq. 44.9. The values of the Langelier index of representative river waters, calculated based on the analyses provided by Suzuki [3] and others, are shown in Figure 44.1.

Charts have been prepared by Powell et al. [7] for calculating the Langelier index of waters varying widely in relevant dissolved species at various temperatures. The historical charts prepared by Powell et al. [7] and the table arranged by Nordell [8] (Table 44.2) are still in use, but correction factors for ionic strength and temperature have been updated based on recent advances in the solubility chemistry of CaCO3 and solution chemistry (e.g., by Schock [9]). While the concept of the Langelier index is correct and helpful, it should be emphasized that a positive value of the index can result from waters of totally different quality. As the pH increases, the Ca2 þ concentration decreases drastically. The corrosion protection characteristics of the resulting CaCO3 film differ accordingly. In other words, waters of different pH, Ca hardness, and MO alkalinity that give the same value of the index have different corrosivity. The buffer capacity and the oxygen concentration of the water and the thickness, composition, and crystalline state of deposits all affect the protectiveness of the deposited CaCO3 film. These facts were pointed out by Stumm [10] as early as 1956 and have not been explained to date either quantitatively or systematically. Instead of using the thermodynamic prediction of the tendency for CaCO3 precipitation in terms of the Langelier index, an experimental method, the marble test, developed by DeMartini [11] and Hoover [12] can be used. The water is

FIGURE 44.1. Values of the Langelier index of representative river waters. Data are primarily from [3].


TABLE 44.2.

593

Data for Rapid Calculation of the Langelier Index Arranged by Nordella A

C

Total Dissolved Solid (ppm)

A

50–300 400–1000

0.1 0.2 B

Temperature ( C)

B

0–1.1 2.2–5.6 6.7–8.9 10.0–13.3 14.4–16,7 17.6–21.1 21.2–26.7 27.8–31.1 32.2–36.7 37.8–43.3 44.4–50.0 51.1–55.6 56.7–63.3 64.4–71.1 72.2–81.1

2.6 2.5 2.4 2.3 2.2 2.1 2.0 1.9 1.8 1.7 1.6 1.5 1.4 1.3 1.2

D

Ca Hardness (ppm as CaCO3)

C

MO Alkalinity (ppm as CaCO3)

D

10–11 12–13 14–17 18–22 23–27 28–34 35–43 44–55 56–69 70–87 88–110 111–138 139–174 175–220 230–270 280–340 350–430 440–550 560–690 700–870 880–1000

0.6 0.7 0.8 0.9 1.0 1.1 1.2 1.3 1.4 1.5 1.6 1.7 1.8 1.9 2.0 2.1 2.2 2.3 2.4 2.5 2.6

10–11 12–13 14–17 18–22 23–27 28–35 36–44 45–55 56–69 70–88 89–110 111–139 140–176 177–220 230–270 280–350 360–440 450–550 560–690 700–880 890–1000

1.0 1.1 1.2 1.3 1.4 1.5 1.6 1.7 1.8 1.9 2.0 2.1 2.2 2.3 2.4 2.5 2.6 2.7 2.8 2.9 3.0

a See [8]. (1) Obtain values of A, B, C and D from the table. (2) pHs ¼ (9.3 þ A þ B) (C þ D). (3) Langelier saturation index ¼ pH pHs.

treated with powdered CaCO3, and, after saturation, changes in pH, alkalinity, and concentration of calcium are measured as an estimate of undersaturation. An improved procedure has been proposed by Merrill and Sanks [13]. Another index used in practice for estimating the corrosivity of water is the stability index or the Ryznar index (RI) developed by Ryznar [14]: RI ¼ 2pHs pH

ð44:10Þ

This index is intended to predict quantitatively the amount of CaCO3 that would be formed and also to predict the corrosivity of waters that are undersaturated with respect to CaCO3. At RI 6.0, CaCO3 is in equilibrium. The deposition of CaCO3 increases proportionately as the index drops < 6, and corrosivity increases as the index rises > 6. Values of 10 or above indicate extreme corrosivity. Both the Langelier and the Ryznar indexes were derived from thermodymamic considerations. The CaCO3 precipitation does not occur unless a certain degree of oversaturation is indicated by these indexes. A new index, the practical saturation index (PSI), has been proposed by Puckorius and

Brooke [15], which uses what they called an equilibrium pH (pHe) in the Ryznar equation [Eq. (44.10)]: PSI ¼ 2 pHs pHe pHe ¼ 1:485 log ðtotal alkalinityÞ þ 4:54

ð44:11Þ ð44:12Þ

The coefficient and the constant in Eq. (44.12) were obtained empirically from the study of hundreds of actual case histories recorded and evaluated by the authors over 12–15 years with concentrated cooling waters, most of which were water treated to a pH level >7.5, often between 8 and 9.5. Because the values of pHe tend to be smaller than those of the measured pH, making the PSI values larger than the Ryznar index values to compensate the oversaturation required for precipitation to occur [cf. Eqs. (44.10) and (44.11)], the CaCO3 precipitation takes place below exactly PSI 6.0, according to the authors. For soft waters, a desired level of the Langelier index can usually be attained by the addition of lime, sodium carbonate, or sodium hydroxide to raise the pH to an appropriate value. For very soft waters, however, the required pH may be too

594


FIGURE 44.2. Relationship between the Langelier index and the corrosion rates of water pipes [18].

high for uses such as tap water unless Ca2 þ ions are added simultaneously. A Langelier index of þ 0.1 þ 0.5 is considered to be satisfactory to provide corrosion protection and at the same time avoid excessive deposition or scaling at elevated temperatures [16]. Data by Flentje [17] and Fujii [18] show that water with a less negative Langelier index is less corrosive. Waters with a slightly negative Langelier index may deposit CaCO3 because of pH fluctuations. The relationship between the Langelier index and the corrosion rates of water pipes obtained by Fujii [18] from a field test at a city waterworks are reproduced in Figure 44.2. Even though the saturation index is positive, the corrosion rate may remain high if the water contains colloidal silica or organic particles, such as algae, because CaCO3 precipitates on them instead of on the steel surface. For waters high in dissolved salt or at high temperature, the CaCO3 film is less protective. The following sections deal with corrosion in soft waters unless otherwise stated. B2.3. pH. The pH of freshwaters is determined by the equilibrium of carbonic species: pH ¼ 6:35 þ

log½HCO3 ð25 CÞ ½CO2aq

ð44:13Þ

where [CO2aq] is the sum of dissolved [CO2] and [H2CO3] in equilibrium with the partial pressure of CO2 gas in the gas phase. The pH of surface waters ranges from 5 to 9, but that of underground waters exposed to a high CO2 partial pressure is often < 6.5 and increases on aeration, which eliminates excessive CO2. Based on the classic data by Whitman et al. [19], it has generally been considered that, within the pH range of

4–10, the corrosion rate is independent of pH, other factors being equal (Fig. 44.3, line a). Regardless of the bulk pH of water within this range, the steel surface is always in contact with an alkaline solution of saturated hydrous ferrous oxide (pH 9.5). Some researchers claimed, however, that water quality (i.e., alkalinity, buffer capacity, and concentrations of Cl and SO42 ) is the primary factor in corrosion at low-flow velocity and that the influence of dissolved oxygen and pH on corrosion rates is secondary. Skold and Larson [20], by polarization resistance measurements, reported the effect of pH on corrosion, as shown by curves b and c in Figure 44.3, where the corrosion rate increases from pH 7 to a maximum at pH 8. The concentrations of NaHCO3 and NaCl were kept constant at 2.5 and 0.5 mmol/L, respectively, throughout the experiment. Carbon dioxide was used to control pH. A similar curve (Fig. 44.3, curve d) was obtained by Pisigan and Singley [21], who conducted immersion tests in a solution containing 1 mmol/L total carbonate species. These authors pointed out that the increase of the corrosion rate in this pH range corresponds to the decrease of the buffer capacity provided by HCO3 . The basis of this conclusion is that HCO3 ions act as an inhibitor (see Section B2.5) Regarding the effect of pH indicated by curves b–d in Figure 44.3, it should be pointed out that the waters used in these experiments contained no calcium, unlike natural waters. Fujii et al. [22] observed an appreciable increase in the corrosion rate with increase of pH from 6.5 to 8.5 (Fig. 44.3, curve e) with soft Tokyo tap water (Ca hardness: 51.8 ppm, MO alkalinity: 39.8 ppm, pH 7.1, Langelier index: 1.4) using CO2 gas or NaOH to control pH. They attributed the increased corrosion rate to the observed enhanced formation


595

FIGURE 44.3. Effect of pH on corrosion of steel in aerated waters. (a) Soft tap water, Cambridge, MA. pH adjusted by additions of HCl and NaOH, 22 C [19]. (b) Aerated water containing 2.5 mmol/L NaHCO3 and 2.5 mmol/L NaOH. pH controlled by continuous addition of CO2 (22–31 C). Linear polarization method applied after 16 days immersion [20]. (c) Aerated water containing 2.5 mmol/L NaHCO3 and 0.5 mmol/L NaOH. pH controlled by continuous addition of CO2 (19–28 C). Linear polarization method applied after 16 days immersion [20]. (d) Water containing NaHCO3 (alkalinity: 100 mg/L as CaCO3). pH adjusted by introduction of CO2 and air. Room temperature, 3 days [21]. (e) Soft tap water, Tokyo, Japan. pH adjusted by additions of CO2 and NaOH. 0.5 in. (12.7 mm OD) pipe, 2 L/min once-through flow. Room temperature, 32 days [22].

and activity of tubercles at higher pH, where corrosion products are more adherent and the buffer capacity is less, resulting in a lower pH at the anodic areas. B2.4. Chloride and Sulfate Ions. As long as diffusion of dissolved oxygen is controlling corrosion, Cl and SO42 ions in concentrations found in freshwaters have essentially no effect on the overall corrosion rate of steel because they do not affect the solubility of oxygen. Localized attack, however, is accelerated as the conductivity of water increases with increase in the concentration of these ions, expanding the effective cathodic areas. Above a critical water velocity, passivation occurs, and the corrosion rate of the passivated steel increases with Cl concentration. For more details, see Section B3. Some authors have shown that, at a given HCO3 concentration in the absence of Ca2 þ , increased Cl or SO42 concentration caused the water to become more corrosive to steel [21, 23, 24]. According to these authors, even in the absence of Ca2 þ , bicarbonates inhibit the corrosion of steel, and Cl and SO42 are detrimental to this inhibitive action. For more details, see Section B2.5.

B2.5. Bicarbonate. Bicarbonate (HCO3 ) is one of the species essential to form a protective CaCO3 film in the presence of Ca2 þ ions at pH above pHs. Some data indicate the effects of HCO 3 ions as an inhibitor of corrosion in the absence of Ca2 þ ions. Although it is not likely that in natural waters an appreciable concentration of HCO3 ions is present without a comparable concentration of Ca2 þ ions, it would be worthwhile to consider this inhibitive action. With solutions that do not contain Ca2 þ , Larson and co-workers [23, 24] reported that at a given concentration of Cl or SO42 the corrosion rate decreased with increase of HCO 3 and at a given HCO3 concentration the corrosion rate increased with increase of Cl or SO42 and leveled off at a certain concentration of these ions. In other words, the 2 / corrosion rate was determined by the Cl /HCO 3 or SO4 HCO3 ratio. These authors, and also Pisigan and Singley [21], suggested that the corrosion product formed in the presence of 2þ was FeCO3, instead of HCO 3 and in the absence of Ca Fe(OH)2, the former being more protective than the latter. The latter authors, having obtained the effect of pH as shown by curve d in Figure 44.3, proposed that the composite effect of

596


pH and alkalinity can be combined into one parameter, the buffer capacity, b which is defined as b ¼ d (alkalinity)/d pH. If the pH is varied at constant alkalinity or total carbonate species, b of the water decreases with pH in the range of typical drinking water systems from a maximum to a minimum at pH corresponding to pK1 and to 12 ðpK1 þ pK2 Þ, respectively, where K1 and K2 are the first and the second dissociation constants of H2CO3. The maximum occurs at pH 6.3 and the minimum at pH 8.3. Curve d in Figure 44.3 shows the highest corrosion rate at pH 8.0, at which b is near its minimum. The beneficial effect of the alkalinity represented by b was explained by its capacity to neutralize the acid generated at the anode by the oxidation of Fe2 þ 2Fe2 þ þ 12 O2 þ 5H2 O ¼ 2FeðOHÞ3 þ 4H þ and to cause additional carbonate precipitation (FeCO3, or CaCO3 in the presence of Ca2 þ ) at the cathode, decreasing the strong driving force for corrosion. B3. Effect of Velocity The effect of water velocity on corrosion of mild steel, from data by Kowaka et al. [25], Fujii et al. [22], Matsushima [26], and Kinoshita et al. [27], is shown in Figures 44.4 and 44.5. According to these data, the critical velocity above which

the corrosion rate decreases because of passivation is in the range of 0.3–0.7 m/s. The critical velocity and the corrosion rate at velocities above critical depend on the concentration of chloride ions that prevent passivation. An example of data by Matsushima [26] is shown in Figure 44.6. As shown by LaQue [28], passivity is not established at any velocity in the presence of high concentrations of Cl , as in seawater. On the other hand, passivation of mild steel occurs at a very low velocity, for example, < 0.1 m/s in high-purity water (e.g., conductivity < 0.5 mS/cm) [29]. Water at very high velocities (e.g., 20 m/s) causes erosion–corrosion by removing corrosion product films through application of shear stress, thereby exposing the steel surface directly to the corrosive environment. It may also remove the metal mechanically, but removal of only corrosion product films is sufficient for erosion–corrosion to proceed. The free access of dissolved oxygen in water increases corrosion several times. The maximum penetration rate approaches the rate that corresponds to the amount of dissolved oxygen supplied from the water without a surface barrier. Erosion–corrosion often occurs at bends or misaligned joints of high-velocity water pipe even if the water does not carry suspended solids. If conditions of high-velocity water are such that low- and high-pressure areas develop and bubbles form and collapse at the metal–water interface (cavitation), the steel is damaged mechanically by the impact pressure exerted by the collapsing bubbles. This type of degradation is called cavitation erosion or cavitation damage. Cavitation erosion can occur purely mechanically but is accelerated by free access of dissolved oxygen. It occurs typically on rotors of pumps and the trailing side of water turbine blades made of steel and cast iron. C. LOCALIZED CORROSION C1. General Localized corrosion of steel in freshwaters occurs in most cases by the action of macrogalvanic cells. The resulting damage tends to be most critical in water pipes because perforation of the pipe wall by localized corrosion (e.g., pitting) means instant leakage of the water carried by the pipe. The major types of localized corrosion of steel in freshwater are pitting under tubercles, groove corrosion of electric resistance welded (ERW) pipes, and pitting at discontinuities in corrosion product films. For other types of localized corrosion in water see Chapter 47.

FIGURE 44.4. Effect of water velocity on corrosion of carbon steel (1). (a) Distilled water þ 10 ppm Cl , 50 C, 14 days [26]. (b) Soft tap water, Tokyo, Japan, room temperature, 67 days [22]. (c) Soft tap water, Amagasaki, Japan, 20 C, 15 days (killed steel) [25]. (d) Soft tap water, Amagasaki, Japan, 20 C, 15 days (rimmed steel) [25].

C2. Pitting under Tubercles Tuberculation or formation of mounds of corrosion products (tubercles) occurs frequently on the steel surface in contact

LOCALIZED CORROSION

597

FIGURE 44.5. Effect of water velocity on corrosion of steel (2). 25 ppm Cl , 30 C, 30 days [27].

with water. Differential aeration cells are then formed. The areas under tubercles, where the supply of dissolved oxygen is limited, undergo localized corrosion by the galvanic action of the surrounding cathodic areas that receive free oxygen supply. Galvanized water pipes corrode similarly after the

zinc coating has been lost, as often occurs within several years of service in soft waters. The penetration rates in millimeters per year (mm/year) and mils per year (mpy) of galvanized city water pipes are listed in Table 44.3 [30]. The periods of protection by the

FIGURE 44.6. Effect of water velocity and chloride concentration on corrosion of carbon steel. Distilled water þ NaCl, 50 C, 14 days [26].

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TABLE 44.3. Penetration Rates of Galvanized Tap Water Pipes Under Tuberclesa

Size (in.) 1 1 1 11/4 1 11/4 11/4 2 1 11/2 5 4 1 6

Maximum Penetration Rate

Wall Thickness (mm)

Maximum Penetration (mm)

Service Period (year)

mm/year

3.2 3.2 3.2 3.5 3.2 3.5 3.5 3.8 3.2 3.5 4.5 4.5 3.2 5.0

3 1 1.5 0.8 Perforated 0.8 1.2 0.8 1.4 1.3 2 Perforated Perforated Perforated

13 5 10 10 14 14 14 14 14 14 16 15 10

0.23 0.20 0.15 0.08 0.23 0.06 0.09 0.06 0.10 0.09 0.13 0.30 0.32

mils/year (mpy) 9 8 6 3 9 2 3 2 4 3 5 12 13

a

From [30].

zinc coating were neglected in calculating the rates. The rates are not high, ranging generally from < 0.1 to 0.3 mm/year (1 mm/year (40 mpy) sometimes occurs without the formation of tubercles in pipes alternately exposed to hot water (or steam) and cold water. Examples are shown in Table 44.4 [39]. The inside surface of the pipe is covered by corrosion product films that are tight, dense, relatively smooth, and dark in color. X-ray analysis identifies g-Fe2O3 and Fe3O4, unlike the usual iron oxides, a- or g-FeOOH, Fe3C4, and large amounts of amorphous substances that form in aqueous media. The degree of crystallization is much higher, as indicated by high intensities of the diffracted beams in X-ray diffraction. The mechanism of this type of corrosion is similar to that of the localized corrosion of mill-scaled steel in seawater (see Chapter 47 Section B4). The surface covered by corrosion product films formed during exposure to hot water or steam tends to act as the cathode to the exposed steel at the defects or discontinuities in the films during exposure to cold water, causing localized corrosion. The corrosion potential of the film-covered areas was measured in the range of 590 to 670 mV versus SCE in 3% NaCl in the laboratory. The localized attack was reproduced in an experiment simulating service conditions [40].

The welded seam of the ERW pipe is heated locally and then cooled rapidly by the ambient air, and its microstructure is much different from that of the parent metal that is not heated during welding. It tends to have more negative, or less noble, potential than the parent metal and, on exposure to corrosive media, such as water and soil, corrodes locally in the form of grooves. This type of localized corrosion does not always occur, but most ERW pipes are potentially susceptible to it A rate in the range of 1–3 mm/year (40–120 mpy) is common and can be as high as 10 mm/year (400 mpy) [35]. The difference of the open-circuit potential is a maximum of 70 mV [35]. The more negative, or less noble, potential of the weld seam has been attributed to the formation of unstable MnS inclusions in the weld by the thermal cycle of the welding operation [36, 37]. Butt-welded pipes that have been heated uniformly in the process of shaping and welding are not susceptible to this type of corrosion. Groove corrosion can be mitigated by minimizing the potential difference between the weld seam and the parent metal. The ERW pipes that are resistant to groove corrosion were developed in the 1970s. The chemical compositions of

C4. Localized Corrosion at Discontinuities in Corrosion Product Films

600


REFERENCES 1. H. Uhlig, D. Triadis, and M. Stern, J. Electrochem. Soc., 102, 59 (1955). 2. S. K. Coburn, in Metals Handbook, Vol. 1, 9th ed., American Society of Metals, Metals Park, OH, 1978, p. 733. 3. T. Suzuki, Private Communication (Database, Kurita Water Industries, Ltd., Tokyo). 4. R. W. Revie and H. H. Uhlig, Corrosion and Corrosion Control, 4th ed., Wiley, Hoboken, N.J., 2008, p. 461. 5. R. A. Pisigan, Jr., and J. E. Singley, J. Am. Water Works Assoc., 77(10), 83 (1985). 6. W. F. Langelier, J. Am, Water Works Assoc., 28, 1500 (1936). 7. S. Powell, H. Bacon, and J. Lill, J. Am. Water Works Assoc., 38, 808 (1945). 8. E. Nordell, Water Treatment, 2nd ed., Reinhold, New York, 1961, p. 287. 9. M. R. Schock, J. Am. Water Works Assoc., 76(8), 72 (1984). 10. W. Stumm, J. Am. Water Works Assoc., 48(3), 300 (1956). 11. F. E. DeMartini, J. Am. Water Works Assoc., 30, 85 (1938). 12. C. P. Hoover, J. Am. Water Works Assoc., 30, 1802 (1938). 13. D. T, Merrill and R. L. Sanks, J. Am. Water Works Assoc., 70 (1), 12 (1978). 14. J. W. Ryznar, J. Am. Water Works Assoc., 36, 472 (1944). 15. P. R. Puckorius and J. M. Brooke, Corrosion, 47, 280 (1991). 16. N. R. Reedy, Mater. Protect. Perform., 12(4), 43 (1973). 17. M. E. Flentje, J. Am. Water Works Assoc., 53, 1461 (1961). 18. T. Fujii, Bosei Kanri (Rust Prevention and Control), 27, 85 (1983). 19. W. Whitman, R. Russell, and V. Altieri, Ind. Eng. Chem., 16, 665 (1924). 20. R. V. Skold and T. E. Larson, Corrosion, 13, 139t (1957). 21. R. A. Pisigan and I. E. Singley, J. Am. Water Works Assoc., 79 (2), 62 (1987). 22. T. Fujii, T, Kodama, and H. Baba, Boshoku Gijutsu (Corrosion Eng.), 31, 637 (1982). 23. T. E. Larson and R. M. King, Corrosion, 10, 110 (1954). 24. T. E. Larson and R. V. Skold, Corrosion, 14, 285t (1958). see also R. V. Skold and T. E. Larson, Corrosion, 13, 139t (1957).

25. M. Kowaka, M. Ayukawa, and H. Nagano, Sumitomo Kinzoku, 21, 185 (1969). 26. I. Matsushima, in Corrosion Protection Handbook (Boshoku Gijutsu Binran), Japan Society of Corrosion Engineering, ed., Nikkan Kogyo Shinbunsha, Tokyo, Japan, 1985. pp. 176, 179. 27. K. Kinoshita, K. Ichikawa, and N. Kitajima, Boshoku Gijutsu (Corrosion Eng.), 32, 31 (1983). 28. F. LaQue, in Corrosion Handbook, H. H. Uhlig (Ed.), Wiley, New York, 1948, p. 391. 29. K. Sakai, S. Morishita, T. Honda, A. Minato, M. Izumiya, K. Osumi, and M. Miki, Boshoku Gijutsu (Corrosion Eng.), 30, 450 (1981). 30. I. Matsushima, paper presented at the Committee on Corrosion and Protection, The Society of Materials Science, Japan, 19, 94 (1980); reproduced in Corrosion Protection Handbook (Boshoku Gijutsu Binran), Japan Society of Corrosion Engineering, Nikkan Kogyo Shinbunsha, Tokyo, Japan, 1985, p. 181. 31. K. Masamura and I. Matsushima, 23rd Annual Symposium on Corrosion and Its Protection, Japan Society of Corrosion Engineering, Tokyo, Nov. 1978, p. 104. 32. R. W. Revie and H. H. Uhlig, Corrosion and Corrosion Control, 4th ed., Wiley, Hoboken, N.J., 2008, p. 276. 33. F. N. Speller, in Corrosion Handbook, H. H. Uhlig (Ed), Wiley, New York, 1948, p. 496. 34. U. R. Evans, The Corrosion and Oxidation of Metals, Edward Arnold, London, 1960, p. 296. 35. K. Masamura and I. Matsushima, Paper No. 75, CORROSION/ 81, NACE, Houston TX, 1981. 36. C. Kato, Y. Otoguro, and K. Kado, Boshoku Gijutsu (Corrosion Eng.), 23, 385 (1974). 37. C. Kato, Y. Otoguro, K. Kado, and Y. Hisamatsu, Corros. Sci., 18, 61 (1978). 38. I. Matsushima, Low-Alloy Corrosion Resistant Steels, Chijin Shokan Co., Tokyo, Japan, 1995, p. 155. 39. I. Matsushima,“A Study on Localized Corrosion of Steel,” Ph.D. Thesis, University of Tokyo, 1982. 40. I. Matsushima, Annual Meeting of Japan Society of Corrosion Engineering, Tokyo, May, 1979, p. 143; reproduced in Localized Corrosion, F. Hine, K. Komai, and K. Yamakawa (Eds.), Elsevier Applied Science, London and New York, 1988, p. 31.

45 CARBON STEEL—CORROSION BY SEAWATER I. MATSUSHIMA* Maebashi Institute of Technology, Maebashi, Japan

A. Introduction B. Corrosion by continuous immersion in seawater B1. Environmental factors B2. Corrosion rate B3. Localized corrosion C. Corrosion of pilings D. Effect of velocity References

A. INTRODUCTION Typical steel structures that are exposed to marine environments are marine piles, offshore structures, vessels, and other structures immersed in seawater. A steel structure in a marine environment may be exposed to five different corrosive zones depending on the position of its parts relative to sea level. The five zones are the atmospheric, splash, tidal, submerged, and seabed-embedded zones. The corrosion characteristics and corrosion behavior of steel are different in each zone. Some kinds of equipment, such as machinery and piping systems that use seawater for industrial purposes, are also subject to seawater corrosion. While well-established countermeasures are available against marine corrosion, such as application of coatings and cathodic protection, corrosion of bare steel in marine environments is important because some steel structures are used without protection and also because corrosion protection systems in marine environments are particularly susceptible to damage and deterioration, leaving considerable exposure time for corrosion before detection and repair.

B. CORROSION BY CONTINUOUS IMMERSION IN SEAWATER B1. Environmental Factors The major variables that affect the corrosion rate of steel are salinity, dissolved oxygen concentration, temperature, pH, carbonate, pollutants, and biological organisms. The characteristics of seawater with respect to these variables are summarized in Table 45.1. Reviews of the variability in seawater at different global locations have been published by Dexter and Culberson [1] and by Dexter [2]. Corrosion of steel in seawater is controlled by the rate of supply of dissolved oxygen to the steel surface, similar to corrosion in freshwaters. The rate of oxygen supply, in turn, is determined by the oxygen concentration in the bulk seawater, the degree of movement of seawater, the diffusion coefficient for oxygen in seawater, and characteristics of corrosion product films on the steel surface as a barrier to oxygen diffusion. The concentration of oxygen in surface waters is usually near the equilibrium saturation value with the atmosphere, which varies inversely with the temperature and salinity of seawater. Because salinity variations in the surface water are relatively small and do not greatly affect oxygen solubility, the temperature is the major factor affecting oxygen concentration. The temperature of the surface waters of oceans varies mainly with the latitude, and the range is from about 2 C in the Arctic to 35 C in the tropics. Accordingly, the equilibrium saturation concentration of oxygen varies from 11 to 6 ppm [1, 2]. The equilibrium concentration of oxygen in water is shown in Table 45.2 as a function of temperature and salinity. The diffusion coefficient for oxygen, on the other hand, increases with increase of temperature, and the corrosion

*Deceased Uhlig’s Corrosion Handbook, Third Edition, Edited by R. Winston Revie Copyright Ó 2011 John Wiley & Sons, Inc.

601

602

CARBON STEEL—CORROSION BY SEAWATER

TABLE 45.1. Environmental Factors in Seawater Corrosiona Salinity

1. Open sea: Variation with horizontal location is small, 32–36 parts per thousand (ppt). 2. Near river outlets: Lower. 3. Variation with depth: Very small.

Dissolved oxygen concentration

1. Surface water: (1) Near the equilibrium saturation concentration with atmospheric oxygen at a given temperature [6 ppm (in the tropics), 11 ppm (in the Arctic)]. (2) Can be supersaturated due to photosynthesis by microscopic plants (up to 200%) and entrainment of air bubbles (up to 10%). 2. Variation with depth: (1) Tends to be undersaturated due to consumption by the biochemical oxidation of organic matter. (2) Goes through a minimum at intermediate depths (400–2400 m deep).

Temperature

1. Surface water: In the open ocean, variations are in the range of 2 to 35 C depending on the latitude, season, currents, and so on. 2. Variation with depth: Drops with depth. The difference with depth and season may be large or small depending on the location.

pH

1. Surface water: (1) Lies between 7.5 and 8.3 in the open ocean depending on the concentration of dissolved CO2 determined by air–sea exchange and photosynthesis activity of plants. (2) Microbiological activity affects the pH value; e.g., lower pHs by the formation of CO2 through the process of biochemical oxidation and higher pH values by the reduction of CO2 through the process of photosynthesis. (3) Affected by pollutants in the coastal waters. 2. Variation with depth: Tends to show a profile similar to that of dissolved oxygen (the biochemical oxidation that consumes dissolved oxygen generates CO2, reducing the pH value).

Carbonate

1. Surface water: Nearly always supersaturated with respect to CaCO3 (200–500%) favored by high pH values and moderate temperatures. 2. Variation with depth: Saturation state with respect to CaCO3 decreases as the result of lower temperature and pH. Undersaturated in deep waters (e.g., below 200–300 m).

Pollutants

1. H2S may be 50 ppm or higher in polluted waters in estuaries, harbors, river mouths, and fitting-out basins. 2. Ammonia may be high in inshore waters and harbors.

Biological organisms

1. Bacteria form bacterial films (slime). 2. Weeds grow from spores. 3. Animals (e.g., barnacles, tube worms, and hydroids) adhere.

a

Prepared based on [2].

rate of steel at a given oxygen concentration approximately doubles for every 30 C rise in temperature [3]. Because of the compensating effects of temperature with respect to the oxygen solubility and the diffusion coefficient, the rate of corrosion of steel in seawater is relatively independent of temperature. The corrosion product film on the steel surface that more or less serves as a barrier to oxygen diffusion and decreases the corrosion rate of steel contains oxides and hydroxides of

iron, and possibly iron sulfide, calcareous deposits, bacterial slime, and macroscopic marine growths. While the film is a physical barrier to oxygen diffusion, oxygen-utilizing bacteria in the film may provide a biochemical barrier by consuming all the oxygen diffusing through the film, resulting in anaerobic conditions at the metal surface. These conditions reduce corrosion by oxygen but may provide a condition where sulfate-reducing bacteria flourish and accelerate corrosion.

TABLE 45.2. Equilibrium Saturation Concentration of Oxygen in Seawater as a Function of Temperature and Salinitya Oxygen Solubility (ppm) at Indicated Salinity (ppt)

Temperature ( C) 0 5 10 15 20 25 30 a

0

8

16

24

31

36

14.6 12.5 10.9 9.5 8.5 7.6 6.8

13.9 11.9 10.4 9.1 8.5 7.2 6.5

13.1 11.3 9.8 8.7 7.7 6.9 6.2

12.4 10.7 9.4 8.2 7.3 6.6 6.0

11.8 10.2 8.9 7.9 7.0 6.3 5.7

11.4 9.9 8.7 7.7 6.8 6.2 5.6

Calculated based on data from [2].

CORROSION BY CONTINUOUS IMMERSION IN SEAWATER

In spite of wide variations of environmental factors affecting the rate of oxygen supply, differences in the corrosion rate of steel under conditions of continuous immersion in seawater at different locations throughout the world are relatively small. LaQue [4] pointed out that this similarity of corrosion rates in seawater everywhere occurs because the controlling factors change in a compensating way. For example, decrease of oxygen concentration by higher water temperatures is compensated by a large diffusion coefficient. High water temperature, which tends to promote higher reaction rates, also enhances the development of protective calcareous deposits and marine growths, which stifle attack [4]. The action of bacteria that grow beneath the original corrosion product film may simply be to take the place of the excluded oxygen in enabling corrosion reactions to continue at the rates commonly observed [4]. Of course, there are some cases where the corrosion rate is exceptionally high or low. For example, the corrosion rate may be very low because of complete coverage of the surface by macroscopic fouling organisms. Significant salinity variations, caused by river discharge or high evaporation, affect oxygen concentration in seawater and hence the overall corrosion rate. Pollutants in inshore waters (e.g., organic material, hydrogen sulfide, and ammonia) affect corrosion by changing the oxygen concentration and/or the nature of corrosion product films as a barrier to oxygen diffusion. Intense photosynthesis by macroscopic marine plants growing in the surface layers of the sea can produce high

603

supersaturation of oxygen in the surface waters in the daylight hours. On the other hand, oxygen concentration of seawater below the surface tends to be lower due to consumption by the biochemical oxidation of organic matter. The vertical profile of oxygen concentration depends on the location but commonly goes through a minimum at intermediate depths. The depth of the oxygen minimum ranges from 400 m in the equatorial eastern Pacific to over 2400 m in the central Pacific and the concentration at the depth of the minimum ranges from 0.16 to 6.4 ppm [2]. The pH of open ocean seawater ranges from 7.5 to 8.3 and has no direct effect on the corrosion rate of steel, but its variation affects the degree of saturation with respect to CaCO3. Spontaneous calcareous deposition is more likely to occur at higher pH values, but the function of the deposits as a barrier to oxygen diffusion is limited because of the presence of organic matter and high salinity of seawater, although enhanced buildup of calcareous deposits under cathodic protection helps to reduce the required protection current, lowering the power consumption. B2. Corrosion Rate Overall corrosion rates of steel continuously immersed in quiescent seawater at many locations throughout the world for periods from < 1 year to 40 years collected from various literature [4–14] are plotted in Figure 45.1. The corrosion rates range from 0.02 to 0.37 mm/year (0.8–14.6 mpy), the average rate being 0.1 mm/year (4 mpy).

FIGURE 45.1. Average corrosion rates of steel continuously immersed in seawater.

604


The average of the corrosion rates in Figure 45.1 in the first 5 years is 0.14 mm/year (5.5 mpy) being very close to 0.125 mm/year (5 mpy), which is the value most commonly used as the expected average rate of corrosion of steel continuously immersed in seawater under natural conditions [4, 15]. The average corrosion rates for periods of > 5 to 10 years, > 10 to 20 years, and > 20 years are 0.07, 0.07, and 0.05 mm/year (2.8, 2.8, and 2.0 mpy), respectively, indicating that the corrosion rate decreases with time. At one time, the corrosion rate of steel in seawater was considered to be linear with time [4]. Larrabee [16] reported in 1962 that the average corrosion rate of submerged sections of 20 H-piles of a steel pier near Santa Barbara, CA, used for offshore oil wells for 23.5 years was 0.038 mm/year (1.5 mpy) and concluded that the corrosion rate of carbon steel in seawater averages 0.05 mm/year (2 mpy) for the first 20 years, then drops to 0.025 mm/year (1 mpy). The decrease of corrosion rate with time has been more clearly shown in corrosion tests in which steel was continuously immersed in seawater at a specific location for different periods of time. Southwell and Alexander [12], who conducted an extensive corrosion test in the Panama Canal Zone, showed that the overall corrosion rates [(average penetration)/(time of exposure)] of structural steel in seawater in 1, 2, 4, 8, and 16 years were 0.14, 0.11, 0.094, 0.084, and 0.077 mm/year (5.5, 4.3, 3.7, 3.3, and 3.0 mpy), respectively (curve b in Fig. 45.1). The corrosion rate was 0.15 mm/year (5.9 mpy) after 1 year of exposure and 0.068 mm/year (2.7 mpy) after 16 years of exposure. From the results of corrosion tests of structural carbon steels in seawater at Wrightsville Beach, NC, Schmitt and Phelps [10] reported that the corrosion rates of structural carbon steel in 1.5, 2.5, 4.5, and 8.5 years were 0.12, 0.10, 0.083, and 0.068 mm/year (4.7, 4.0, 3.3, and 2.7 mpy), respectively (curve c in Fig. 45.1). B3. Localized Corrosion One of the common causes of pitting of carbon steel is differential aeration cells formed by nonuniformity of corrosion product films or spotty coverage of fouling organisms on the surface. Mill scale, if it has not been removed prior to exposure to seawater, accelerates corrosion at bare spots or breaks in the scale through galvanic action, resulting in pits. Galvanic action tends to decrease with time because of the development of insulating calcareous deposits as a result of cathodic reaction at the scaled surface [4] and also because mill scale is undermined by corrosion products that cause it to be spalled off. For this reason, the effect of mill scale is more pronounced during the early stages of exposure. From a literature survey, Fink [15] concluded that the pitting attack on bare steel is frequently 0.25–0.38 mm/ year (10–15 mpy) and that the presence of mill scale significantly increases the rate of pitting, the penetration rate being

0.5 mm/year (20 mpy). Assuming the average corrosion rate of 0.125 mm/year (5 mpy), the pitting factor (the ratio of deepest metal penetration to average metal penetration) is 2–3 for bare steel and 4 for mill-scaled steel. According to LaQue [4], the pitting rate decreases with time, and a normal pitting factor for exposure of 10 years would be 2.5 for descaled steel and 3.5 for steel exposed with mill scale. He also stated that, as a general rule, steel exposed with mill scale will be pitted about three times as deeply as descaled steel in a short period of exposure and that this ratio decreases as the exposure is prolonged to become 1.5–1 for a 10-year-exposure period. The values of the pitting factor found in 15-year tests in seawater at Halifax (N.S., Canada), Auckland (New Zealand), Plymouth (England), and Colombo (Sri Lanka) by the Committee of the Institution of Civil Engineers (London) [17] were 2.0–2.9 for pickled specimens and 2.1–5.6 for specimens with mill scale, in fair agreement with the values indicated by Fink and LaQue. The pitting factor based on the depth of the deepest pit and the average of the 20 deepest pits found by Southwell et al. [11, 12] in the tropical sea of the Panama Canal Zone for machined and mill-scaled structural steel are summarized in Table 45.3. In these data, the pitting factors are much greater than those summarized by Fink and LaQue.

C. CORROSION OF PILINGS The corrosion rates in the region between high and low tide, which is alternately immersed and not immersed (the tidal zone), and in the area above it which receives seawater spray (the splash zone) are very high because of constant wetting. The corrosion attack in the tidal zone occurs primarily during the periods of atmospheric exposure because of the presence of the thin seawater film and an abundant supply of oxygen. The splash zone is under such conditions all the time. However, bare (uncoated) vertical members, such as marine piles, that extend from the submerged zone through the tidal zone into or beyond the splash zone frequently show characteristic corrosion behavior as first reported by TABLE 45.3. Pitting Factor of Structural Steel in Seawater in the Panama Canal Zonea Exposure (year) Surface Finish

1

2

4

8

16

Machined Mill scaled

12 25

17 14

8 10

6 11

3b –c

a

Based on data from [11, 12]. Pickled specimen. Pit perforated in machined specimen. The pitting factor for cast steel, which corroded more or less the same as carbon steel, is 3. c Perforated. b

CORROSION OF PILINGS

605

FIGURE 45.2. Corrosion profile of steel piling in seawater. (a) Uncoated piling. (Reproduced with permission from [5]. Copyright Ó NACE International.) (b) At coating defect of piling coated with an electrically insulating substance.

Humble [5] in 1949. Relative corrosion rates of a bare marine pile exposed to the atmospheric, splash, tidal, submerged, and mud (seabed) zones are shown by line a in Figure 45.2. Of these zones, corrosion at the splash zone is the severest, as expected, but corrosion in the tidal zone is relatively mild. It is generally believed that the low corrosion rate in the tidal zone is the result of galvanic protection by the part near the top of the submerged zone that shows a corrosion peak [5]. The part in the tidal zone acts as the cathode probably because the corrosion products (iron oxides) are oxidized to higher oxidation states during the periods of exposure to the atmosphere, resulting in a more noble corrosion potential. Then, when the part in the tidal zone is submerged during the periods of high tide, it acts as the cathode, with the reduction of the oxides on its surface [18]. If the pile is coated with electrically insulating substances, such as organic coatings, the exposed steel at coating defects in the tidal zone corrodes just as severely as in the splash zone [line b in Fig. 45.2] [19]. High corrosion rates near the top of the submerged zone are not always observed. Zen [14] investigated corrosion of piles used at 43 ports throughout Japan for times of up to 40 years and classified the vertical corrosion profile of bare steel marine pilings (sheet and pipe piles) below the low water level into five types. These types may be rearranged to the following three types:

Type 1 Sharp corrosion peak within several tenths of a meter below the low water level (corrosion rate 0.5–1.0 mm/year, 20–40 mpy) and mild corrosion below this level (corrosion rate < 0.1 mm/year, < 4 mpy) Type 2 Corrosion rate highest just below the low water level (0.1–0.2 mm/year, 4–8 mpy), gradually decreasing with depth Type 3 Low and almost flat corrosion rate below the low water level (< 0.1 mm/year, < 4 mpy) Type 1, corresponding to the corrosion profile of line a in Figure 45.2, usually applied to pilings in shallow water, < 3 m deep, with the lowest level of concrete coverage staying above the mean water level. When the pilings were covered with concrete below this level or the water depth was > 5 m, the sharp corrosion peak at the top of the submerged zone was not found, resulting in the corrosion profile of type 3. At locations where the difference in water level of high and low tides was small, the sharp corrosion peak did not occur. The presence of a freshwater layer at the water surface caused by river discharge, on the other hand, caused the corrosion peak to occur even in pilings in deep water. The corrosion profile of type 2 was intermediate between types 1 and 3.

606


TABLE 45.4. Corrosion Rates of Carbon Steel Pilings and Test Specimens in Different Corrosion Zonesa,b Vertical Position

Material

Number of Data

Atmospheric zone

Coupon Average

19 19

0.4–16 —

0.128 0.128

Splash zone

Sheet pile Pipe pile H-pile Coupon Average

8 1 2 16 27

6–40 8 5–7 0.4–15 —

0.112 0.25 0.198 0.363 0.272

Tidal zone


35 4 2 27 68

5–40 3–8.5 5–7 0.4–16 —

0.044 0.070 0.055 0.137 0.083

Low water level

Sheet pile Average

42 42

5–42 —

0.047 0.047

Immersed zone


59 5 3 61 128

5–42 3–8.5 5–23.6 0.3–16 —

0.039 0.062 0.049 0.143 0.090

Mud zone

H-pile Coupon Average

2 3 5

Time of Exposure (years)

Corrosion Rate (mm/year)

5–7 3–5 —

a

From [13]. Averages of corrosion rates compiled from literature.

b

FIGURE 45.3. Effect of velocity on corrosion of steel and cast iron in seawater. (a) Carbon steel tested for 36 days at 23 C [4]. (b) Carbon steel tested for 30 days at ambient temperature [20]. (c) Cast iron tested for 7 days at 25 C [21].

0.033 0.103 0.075

REFERENCES

Corrosion rates of carbon steel pilings and test specimens in different corrosion zones, compiled by a committee of The Kozai Club (organization composed of major steel manufacturers and traders in Japan) [13] from the literature, are listed in Table 45.4. Typical corrosion rates of bare pilings are 0.1–0.2 mm/year (4–8 mpy) in the atmospheric zone, 0.3–0.5 mm/year (12–20 mpy) in the splash zone, 0.1 mm/year (4 mpy) in the tidal zone, 0.5–1 mm/year (20–40 mpy) at the top of the submerged zone (down to 1 m deep) when the corrosion peak appears [otherwise 0.1 mm/year (4 mpy)], 0.1 mm/year (4 mpy) in the submerged zone < 1 m, and 0.05 mm/year (2 mpy) in the mud zone. The pitting factor in the tidal and splash zones is in the range of 2–3.

D. EFFECT OF VELOCITY The corrosion rate of steel by seawater increases with increase of velocity until a critical velocity is reached, beyond which there is little further increase in corrosion rate. The trend of the velocity effect is shown in Figure 45.3 [4, 20, 21]. This behavior is different from that in freshwaters, where corrosion decreases above a critical velocity by passivation (see Sections A and B3, in Chapter 44). In seawater, passivity is not established at any velocity because of the high concentration of Cl . The limiting rate of corrosion corresponds to the maximum rate of oxygen supply by diffusion. When the velocity increases above a critical value, corrosion of steel markedly increases because of erosion–corrosion, even in the absence of solid particles. The critical velocity depends on the state of flow, but it is said to be 20 m/s [20]. At areas where flow is disturbed (e.g., bends and joints in piping), erosion–corrosion may occur at much lower velocities. The expected maximum corrosion rate of carbon steel piping under velocities up to 4 m/s is 1 mm/year (40 mpy), but above this velocity, erosion– corrosion would occur at areas of flow disturbance.

607

REFERENCES 1. S. C. Dexter and C. H. Culberson, Mater. Perform., 19(9), 16 (1980). 2. S. C. Dexter, in ASM Handbook, Vol. 13, American Society for Metals, Metals Park, OH, 1987, p. 893. 3. G. Skaperdas and H. Uhlig, Ind. Eng. Chem., 34, 748 (1942). 4. F. L. LaQue, in Corrosion Handbook, H. H. Uhlig (Ed.), Wiley, New York, 1948, p. 383. 5. H. A. Humble, Corrosion, 5, 292 (1949). 6. J. C. Hudson, J. Iron Steel Inst., 166, 123 (1950). 7. J. C. Hudson and J. F. Stanners, J. Iron Steel Inst., 180, 271 (1955). 8. C. P. Larrabee, Proc. ASTM, 44, 161 (1944). 9. C. P. Larrabee, Corrosion, 14, 501t (1958). 10. R. J. Schmitt and E. H. Phelps, J. Metals, 22(3), 47 (1970). 11. C. R. Southwell, B. W. Forgeson, and A. L. Alexander, Corrosion, 16, 512t (1960). 12. C. R. Southwell and A. L. Alexander, Mater. Prot., 9(1), 179 (1970). 13. Committee on Corrosion of Marine Structures, Kozai Club, Tokyo, Japan, Report on Corrosion of Marine Structures, 1974. 14. K. Zen, Submerged Marine Structures—Corrosion Survey and Corrosion Protection (in Japanese), Kajima Shuppan Kai, Tokyo, Japan, 1974. Also see Boshoku Gijutsu (Corros. Eng.), 17, 103 (1968), 18,194 (1969), 20, 414 (1971), 20, 453 (1971), 21, 466 (1972), 22, 55 (1973), 22, 428 (1973). 15. F. W. Fink, Corrosion of Metals in Sea Water, PB 171344, Battelle Memorial Institute, Columbus, OH, 1960. 16. C. P. Larrabee, Meter. Prot., 1(12), 95 (1962). 17. Deterioration of Structures in Sea Water, 18th Report of the Committee of the Institution of Civil Engineers, London, 1938; cited in [4]. 18. I. Matsushima, Low-Alloy Corrosion Resistant Steels, Chijin Shokan Co., Tokyo, Japan, 1995, p. 122. 19. C. P. Larrabee and F. L. LaQue, the discussion to [5]. 20. K. Ichikawa, K. Nagano, S. Kobayashi, and N. Kitajima, Ebara Eng. Rev., No. 85, 2 (1973). 21. M. Miyasaka and N. Takahashi, private communication.

46 CARBON STEEL—CORROSION BY SOILS I. MATSUSHIMA* Maebashi Institute of Technology, Maebashi, Japan

A. Introduction B. Disturbed soils B1. Overall corrosion B2. Localized corrosion C. Undisturbed soils References

different parts of a structure, may increase the rate of penetration by an order of magnitude.

B. DISTURBED SOILS B1. Overall Corrosion

A. INTRODUCTION Typical steel structures used underground in contact with the soil are pipelines, utility pipings, and pilings. While most pipelines are protected against corrosion with coatings and cathodic protection, driven pilings are commonly used in the bare condition. Corrosion rates of steel in soils vary to a marked degree with the kind or type of soil, being affected by many environmental factors such as soil composition, pH, moisture, and so on. Corrosion proceeds basically by the action of water and oxygen, as in water and in the atmosphere, but localized corrosion or pitting is more likely to occur because of nonhomogeneity of the surrounding soils and nonuniform contact of the metal with soils. These effects are more pronounced in disturbed soils than in undisturbed ones, so that pitting in buried pipes is more severe than in pilings driven into undisturbed natural soils. Soils are not severely corrosive environments, and the overall corrosion rates of buried steel are normally far less than 0.1 mm/year (4 mpy). However, localized corrosion caused by macrogalvanic cells, due to differential aeration and other causes that develop potential differences between

Major factors that govern corrosivity of a given soil are porosity (aeration), electrical conductivity or resistivity, dissolved salts, moisture, and acidity, or alkalinity [1]. There are mutual relationships among these factors (e.g., a porous soil may retain more moisture and a soil with a high dissolved salt content has a high conductivity). The same factor may accelerate or retard corrosion, for example, a porous and hence well-aerated and moist soil tends to increase the initial corrosion, but the corrosion product films formed in a well-aerated soil may be more protective than those in an unaerated soil, reducing corrosion, particularly pitting, in the long term. An extensive series of field tests on various metals was conducted by the National Bureau of Standards (NBS; now National Institute of Standards & Technology, NIST) starting in 1910 at many locations in the United States and included almost all types of soils. The results of the tests [2] on 6-in.(152-mm-) long, 3-in.-(76-mm-) diameter open-hearth steel pipe tested for 12 years from 1922 at 44 locations are summarized in Table 46.1 and Figure 46.1. The overall corrosion rates range from 0.003 to 0.063 mm/year (0.1 to 2.5 mpy), the average being 0.02 mm/year (0.8 mpy). With respect to the overall corrosion rates, corrosion by soils is relatively mild. The data on the effects of environmental factors, that is, soil resistivity, internal drainage, and air–pore space, shown

*Deceased. Uhlig’s Corrosion Handbook, Third Edition, Edited by R. Winston Revie Copyright 2011 John Wiley & Sons, Inc.

609

610

CARBON STEEL—CORROSION BY SOILS

TABLE 46.1. Summary of NBS Field Test Results on Open-Hearth Steel Tested for 12 Years at 44 Locations in the United Statesa Overall Corrosion Rate (mm/year) Maximum

0.063

Minimum

0.003

Average

0.020

Merced silt loam, Buttonwillow, CA. Everett gravelly sandy loam, Seattle, WA 44 locations

Pitting Rate (mm/year) Maximum Minimum

>0.45b 0.033

Average

0.143

Muck, New Orleans, LA Everett gravelly sandy loam, Seattle, WA 44 locations

a

Original data based on [2]. Perforated.

b

in Table 46.2 are based on the NBS tests. No single factor controls the overall corrosion rate. From resistivity measurements of soils of known corrosivity, it was found quantitatively that corrosivity is higher for soils of lower resistivity [3, 4], suggesting that soil resistivity can be used as a rough index of corrosivity of soils. It should be noted, however, that the corrosivity was based on field experience with oil and gas pipelines [3] or gas- and waterdistributing systems [4], where localized corrosion was supposedly the main concern. As indicated in Table 46.2, soil resistivity is not a criterion of corrosivity of soils with respect to overall corrosion.

Tables to estimate corrosivity of soils are given in a German Industrial Standard (DIN 50929, Teil 3). One of the tables lists rating indices for each of 12 factors, and by adding the relevant indices and by referring to other tables, the corrosion probability of a given soil can be estimated. The 12 factors include kind of soil (e.g., clay content), resistivity, water content, pH, acidity and alkalinity, sulfide content, and Cl and SO42 contents. B2. Localized Corrosion Localized corrosion, usually in the form of pits, tends to be severe in disturbed soils, and particularly in pipelines, which pass through soils that are different from place to place along the route, thereby enhancing the possibility of the establishment of macrogalvanic cells. Major causes of the formation of macrogalvanic cells in undergound pipelines are illustrated in Figure 46.2. Pitting, rather than general reduction of thickness, occurs in the anodic areas because corrosion current tends to leave the anodic areas at discrete points of low pipe-to-soil resistance caused by nonuniform contact of the metal with the soil or low resistivity of the soil at localized areas (see Chapter 47, Section B5). As shown in Table 46.1 and Figure 46.1, the specimens used in the NBS tests were severely pitted, the highest penetration rate being >0.45 mm/year (>18 mpy). The pitting factors ranged from 3.0 to 22.4 (average: 8.5; 80% of the data points are between 5 and 10), which are much higher than in waters. These high pitting factors were apparently caused by the formation of differential aeration

FIGURE 46.1. Overall corrosion rates and pitting rates of open-hearth steel exposed to 44 soils for 12 years. Original data are based on [2].

DISTURBED SOILS

TABLE 46.2.

Effects of Environmental Factors on Corrosion of Steel in Soilsa Overall Corrosion Rate (mm/year)

Environmental Factor Resistivity (Wcm) 12000 Drainage Very poor Poor Fair Good Air–pore space (%) 30 a b

611

Maximum Pitting Rate (mm/year)

Maximum

Minimum

Average

Maximum

Minimum

Average

0.063 0.058 0.033 0.036

0.018 0.006 0.005 0.003

0.033 0.017 0.018 0.014

0.31 >0.45b 0.23 0.26

0.11 0.05 0.06 0.03

0.20 0.14 0.14 0.11

0.058 0.037 0.063 0.022

0.038 0.010 0.018 0.003

0.046 0.024 0.022 0.010

>0.45b 0.23 0.31 0.18

0.16 0.05 0.08 0.03

0.28 0.14 0.16 0.11

0.033 0.063 0.037 0.058 0.038

0.010 0.009 0.006 0.012 0.004

0.021 0.024 0.017 0.025 0.013

0.20 0.31 0.26 >0.45b 0.23

0.05 0.10 0.05 0.10 0.03

0.13 0.17 0.15 0.20 0.09

Original data are based on NBS field tests [2] on open-hearth steel for 12 years at 44 locations in the United States. Perforated.

cells, as suggested by Romanoff [5]. Pitting by differential aeration in operating pipelines would be more severe than in the tests because larger cathodic areas would be available than in the 6-in.-long specimens used in the tests. The rate at which pits grow in the soil under a given set of conditions tends to decrease with time, as reported by Romanoff [2]. The pit depth–time curves conform to the equation P ¼ ktn , where P is the depth of the deepest pit at time t and k and n are constants that depend on the characteristics of the soil. Examples of the pit depth–time curves obtained in the NBS tests for ferrous metals are reproduced in Figure 46.3 [2]. The most damaging macrogalvanic corrosion cells are formed where there is contact between a pipeline and one or

more foreign metallic structures when the pipeline and the structures are not isolated from each other (case A in Fig. 46.2). The damage is catastrophic when buried piping is in contact with reinforcing bars (rebars) in concrete foundations or walls. Being in the alkaline environment of concrete, rebars are passivated and thus have a noble potential constituting a cathode of very large surface area. The penetration rates observed in buried utility piping are shown in Figure 46.4 [6]. A rate of 1 mm/year (40 mpy) is common and the rate reaches over 3 mm/year (120 mpy), being an order of magnitude larger than those found in the NBS tests (see Fig. 46.1). Pipelines may suffer bacterial corrosion and stray current electrolysis, which are treated elsewhere in this book.

FIGURE 46.2. Major causes of the formation of macrogalvanic cells in underground pipelines.

612

CARBON STEEL—CORROSION BY SOILS

50

Number of cases

40

30

20

10

0

1

2

3

Max. penetration rate, mm/year

FIGURE 46.4. Maximum pitting rates of buried utility pipings [6].

values [7–10]. The highest rate of penetration for the zone above the water table is 0.37 mm/year (15 mpy), whereas below the water table, the rate is only 0.12 mm/year (4.7 mpy). FIGURE 46.3. Pit depth–time curves for ferrous metals exposed to soils of different aeration [2]. (Reproduced courtesy of National Institute of Standards and Technology, Technology Administration, U.S. Department of Commerce.)

C. UNDISTURBED SOILS Undisturbed soil is the environment of steel piles driven below the water table. Factors affecting the corrosivity of undisturbed soils are basically the same as those of disturbed soils. When a pile is isolated from subsurface structures, corrosion tends to be low regardless of the nature of the soil because of low availability of oxygen. Corrosion by differential aeration can occur by the effect of a disturbed and aerated zone above the groundwater level. Coupling with subsurface structures also establishes a galvanic system. Piles at the undisturbed soil zone, however, do not corrode severely because the corrosion current is spread over large surface areas. The pitting tendency is less than in disturbed soils because of the uniform metal-to-soil contact. It is difficult or impossible to expose existing piling for examination, and as a result, corrosion data are scarce. Figure 46.5 has been compiled from limited reported

FIGURE 46.5. Maximum penetration of driven steel pilings. Compiled from [7–10].

REFERENCES

Note: A source list of 1172 English language journal articles on underground corrosion of steel and other metals as well as its causes and prevention was published in 10 parts by Williams [11].

REFERENCES 1. R. W. Revie and H. H. Uhlig, Corrosion and Corrosion Control, 4th ed., Wiley, Hoboken, N.J., 2008, p. 206. 2. M. Romanoff, Underground Corrosion, National Bureau of Standards Circular 579, National Bureau of Standards, Gaithersburg, MD, 1957. See also: Richard E. Ricker, Analysis of Pipeline Steel Corrosion Data from NBS (NIST) Studies Conducted between 1922 – 1940 and Relevance to Pipeline Management, NISTIR 7415, National Institute of Standards and Technology, U.S. Department of Commerce, Gaithersburg, MD, May 2, 2007; and J. Res. Natl. Inst. Stand. Technol., 115(5), 373 (September–October 2010). 3. F. O. Waters, Corrosion, 8, 407 (1952). 4. G. H. Booth, A. W. Cooper, P. M. Cooper, and D. S. Wakerley, Br. Corros. J., 2, 104 (1967).

613

5. M. Romanoff, Underground Corrosion, National Bureau of Standards Circular 579, National Bureau of Standards, Gaithersburg, MD, 1957, p. 11. 6. I. Matsushima, Koatsu Gas, 19, 557 (1982); cited in Corrosion and Corrosion Protection Databook, I. Matsushima (Ed.), Japan Society of Corrosion Engineering, Tokyo, Japan, 1995, p. 143. 7. M. Romanoff, Corrosion of Steel Pilings in Soils, National Bureau of Standards Monograph 58, National Bureau of Standards, Gaitnersburg, MD, 1962. 8. Tetudo-Denka-Kyokai, A Study on Long-term Corrosion of Steel by Soils, 10th Report, Tokyo, Japan, 1980, p. 11; cited in Corrosion and Corrosion Protection Databook, I. Matsushima (Ed.), Japan Society of Corrosion Engineering, Tokyo, Japan, 1995, p. 31. 9. Y. Ohsaki, Corrosion of Steel Pilings, Kozai Club, Tokyo, Japan, 1980. 10. Y. Mori, Doboku-Gijutsu-Shiryo, 23(6), 33 (1981); cited in Corrosion and Corrosion Protection Databook, I. Matsushima (Ed.), Japan Society of Corrosion Engineering, Tokyo, Japan, 1995, p. 33. 11. J. Williams, Mater. Perform., 21(1), 40; (2), 27; (3), 23; (4), 25; (5), 23; (6), 9; (7), 40; (8), 37; (9), 39; (10), 52 (1982).

47 LOCALIZED CORROSION OF IRON AND STEEL I. MATSUSHIMA* Maebashi Institute of Technology, Maebashi, Japan

A. General characteristics B. Macrogalvanic cells B1. Types of macrogalvanic cells B2. Bimetallic contact B3. Localized corrosion at welded joints B4. Discontinuous surface films B5. Differential aeration B6. Differential pH C. Localized Corrosion by Other Causes C1. Other types of macrogalvanic cells C2. Nongalvanic types D. Maximum Possible Penetration Rate References

In certain situations, the anodic area of macrogalvanic cells corrodes more or less evenly, resulting in localized corrosion, but in some other situations, specific sites in an anodic area corrode preferentially due to inhomogeneity of the metal surface and/or the environment to which the anodic area is exposed. For example, the anodic areas of differential aeration cells in steel pipe buried in the ground usually form discrete pits at localized sites where the metal/soil contact resistance is lower than at the rest of the anodic areas. Localized corrosion of iron and steel may occur by the severe corrosive action of the environment concentrated at specific sites without the formation of macrogalvanic cells. The formation of deep pits in cylinder liners (coolant side) of diesel engines by cavitation–erosion under corrosive action is a typical example [1].

B. MACROGALVANIC CELLS A. GENERAL CHARACTERISTICS B1. Types of Macrogalvanic Cells Localized corrosion of iron and steel, typically pitting, usually occurs by the action of macrogalvanic cells. In a macrogalvanic cell, the anodic and cathodic areas are macroscopic, and their locations are fixed, whereas in a microgalvanic cell the anodic and cathodic sites are microscopic and their locations change randomly with time. Relatively small fixed anodic areas surrounded by or connected to relatively large cathodic areas undergo corrosion. Unlike pitting of passive metals, such as stainless steels, where localized breakdown of passivity and the resulting formation of passive–active cells of large potential difference cause deep pits, localized corrosion of iron and steel tends to be shallow in most cases, as the potential difference of various macrogalvanic cells is not as large.

The macrogalvanic cells that frequently cause localized corrosion in steel in service are commonly formed by bimetallic contact, an inhomogeneous steel matrix (typically at welded joints), discontinuous surface films, differential aeration, and differential pH caused by combinations of an alkaline (pH > 10) and a near-neutral (pH in the range of 5–9) environment that leads to the formation of passive– active cells. An alkaline environment is typically provided by mortar or concrete. B2. Bimetallic Contact Bimetallic corrosion of steel that occurs when the steel is coupled to a more noble metal usually takes the form of


615

616

LOCALIZED CORROSION OF IRON AND STEEL

general attack over the steel surface. If the steel in bimetallic contact with a noble metal is coated with organic materials, such as paint, localized corrosion, commonly in the form of severe pitting, occurs at holidays (i.e., defects) in the coating because of the large cathode–anode area ratio. In one case, 350-mm-diameter sewage piping of 7.9 mm wall thickness, internally coated with coal tar epoxy resin, was perforated in 7 months, and the penetration rate was as high as 13.5 mm/year (530 mdd) [2]. The internal surface of the piping at each butt-welded joint was lined with 400-mmwide type 304 stainless steel instead of the resin paint coating that would be damaged by welding. The localized attack apparently proceeded at a coating holiday. Thus, the cathode–anode area ratio was extremely large. In addition, the specific resistance of the sewage was low, 60 W cm, allowing the flow of large galvanic current. Bimetallic contact tends to accelerate localized attack caused by some other types of macrogalvanic cells. For example, localized corrosion at welds proceeds more rapidly when a more noble metal is coupled, because the resulting shift of the electrode potential of the anodic welds in the noble direction causes a higher corrosion rate of the welds and much deeper localized attack. B3. Localized Corrosion at Welded Joints When the weld metal and/or the heat-affected zone (HAZ) are less noble than the steel (parent metal), the former corrodes selectively. Since both are ferrous metals, the difference in the open-circuit potential between the anode and the cathode would not be large, most often a few hundredths of a volt, but the penetration rate could be high because a small shift of the potential of the anode in the noble direction causes a large increase of the dissolution rate due to the small Tafel slope of the anodic polarization curve for steel (see Section D in this chapter). As mentioned in Section C3 in Chapter 44, the welded seam of electric resistance welded (ERW) pipe is susceptible to groove corrosion. With a potential difference of 30 mV, the maximum penetration rate could be as high as 10 mm/ year (400 mpy), a rate of 1–3 mm/year (40–120 mpy) being quite common. In ERW pipes developed to resist groove corrosion, the potential difference is practically nil [3]. In welded joints, the weld metal and/or HAZ corrodes selectively if it is less noble than the parent metal, which is commonly much larger in surface area than the weld metal and HAZ. The corrosion potentials of the welding materials are almost the same as those of the parent metals but frequently are not exactly the same. Usitalo [4] reported a maximum of 40 mV less noble potential for the welding materials than for the steel plates for shipbuilding. The potential of steel tends to shift in the active direction as a result of the thermal cycle during welding operations, making the HAZ less noble than the parent metal.

The selective attack of weld metal can be prevented by adjusting the weld metal composition so that the potential is slightly positive to that of the parent metal, but not too positive, so as to avoid a detrimental effect on the HAZ. To minimize the potential drop at the HAZ, the chemical composition of the steel plate must be appropriate. R€as€anen and Relander [5] made the welding material for shipbuilding steel [tensile strength: 490 N/mm2 (50 kgf/mm2)] slightly more noble than the parent metal by adding Cu and Ni, which make the potential more noble, and decreasing Si, which makes the potential less noble. The chemical composition of the steel plate was also adjusted by specifying the maximum amount of Cu and the minimum amount of Si to make the steel slightly negative to the weld metal. At the same time, the Si, Mn, and S contents were limited and some Ce was added to minimize the potential drop at the HAZ. Some V and Nb were added to counteract the decrease of tensile strength caused by the low Mn content. Chemical Composition (%) Cu

Ni

Si

Mn

S

Others

0.6 0.4 0.25 0.8 Welding material (example) Steel plate 0.3/0.5 — 0.2/0.3 51.1 3 mm/year (120 mpy) [13]. If limited areas of reinforcing bars in concrete structures are neutralized by the action of water penetrating through cracks in the concrete, localized corrosion results at such areas. Along the same lines, piping in constantly wet concrete floors of kitchens in restaurants and in the concrete bottoms of swimming pools may be perforated due to local neutralization by penetrating water. If piping installed in concrete emerges into soil or water, corrosion may occur in the latter environments. The localized attack tends to be concentrated and severe if the piping goes through a thin water layer just outside the concrete, as in the case of a wet concrete floor.

C. LOCALIZED CORROSION BY OTHER CAUSES C1. Other Types of Macrogalvanic Cells The presence of chlorides in alkaline environments causes pitting corrosion in passivated steel by locally breaking down passivity. The mechanism is similar to that of pitting corrosion of stainless steels in neutral chloride environments. Localized damage of reinforcing bars in concrete contaminated by chlorides (e.g., use of sea sand in concrete without sufficient washing) and pitting in steel chemical equipment handling crude caustic soda containing

chlorides are examples. Pitting corrosion occurs in water (with or without chlorides) inhibited by passivating-type inhibitors (e.g., chromates and nitrites) at insufficient concentrations. Below a certain critical concentration, passivators behave as active depolarizers and increase the corrosion rate at localized areas. Steel sheets coated with metals more noble than steel (e.g., nickel, silver, copper, lead, or chromium) are attacked at exposed pores by the galvanic effect of the noble coatings. Similar accelerated corrosion occurs in steel coated with thin paint films (e.g., < 100 mm thick) at coating holidays. Steel coated with thin paint films having some degree of ionic conductance exhibits noble electrode potentials whether or not the paint is pigmented with antirusting compounds of the passivator type. The noble potential may be associated with passivity, but its mechanism, particularly when passivating pigments are not used, has not yet been clarified. C2. Nongalvanic Types Localized corrosion may occur by the locally concentrated direct action of corrosive environments. The bottoms of pipes or vessels in contact with gasoline containing suspended water are corroded locally by the sedimentary water. The high corrosivity of this water is caused by the good supply of dissolved oxygen from the gasoline, in which the solubility of oxygen is as high as six times that in water. Steam return lines are pitted by deposited droplets of condensate at locations where condensation starts. Steam condensate is corrosive if it contains CO2, which is generated in the boiler and contained in the steam. The penetration rate may reach a few millimeters a year [14]. On reaching lower temperatures, steam condensate accumulates at the bottom of the horizontal line to cause elongated general thinning of the bottom areas [14]. Steam lines are pitted similarly if condensate is generated by lowering of temperature. Heat exchanger tubes carrying steam to heat fluid outside the tubing may suffer the same damage. Stray-current corrosion or electrolysis of underground pipelines commonly takes the form of pitting because the discharging current from the pipeline chooses locally distributed paths of low pipe-to-soil resistance. Erosion–corrosion causes localized attack at areas where corrosion product films are removed allowing easier access by corrosive species from the environment. Examples are seen at the bottoms of slurry lines, at bends of piping carrying high-velocity water, and at the inlets of heat exchanger tubes handling corrosive fluids. Cavitation–erosion that occurs typically on diesel engine cylinder liners (cooling water side), on rotors of pumps, and on the trailing side of water turbine blades is accompanied by numerous deep pits. Fretting corrosion caused by oscillatory motion is characterized by the formation of pits.

REFERENCES

FIGURE 47.1. Steady-state anodic current density as a function of the potential of mild steel in deaerated 3% NaCl at 25 and 60 C [10].

D. MAXIMUM POSSIBLE PENETRATION RATE

TABLE 47.1. Highest Penetration Rates for Various Types of Localized Corrosion in Mild Steel Observed in Practice

Type of Localized Corrosion Localized corrosion under tubercles in tap water pipe Localized corrosion under tubercles in hot water pipe Localized corrosion in seawater Localized corrosion by the galvanic action of oxides in steam/industrial water pipe Localized corrosion of underground service pipe Groove corrosion at the weld of electric resistance welded water pipe

The maximum penetration rate at 25 C predicted by Figure 47.1 is 1 mm/year (40 mpy) at 650 mV and 4 mm/year (160 mpy) at 600 mV (vs. SCE). The actual penetration rates of localized corrosion caused by macrogalvanic cells would be lower than the predicted maximum rates because an incubation period usually exists, during which galvanic cells are formed, or because the anodic area is not always fully active. Unfortunately, data on the corrosion potential during service are not always available. The highest penetration rates observed in some types of localized corrosion mentioned earlier are 0.3–10 mm/year (12–400 mpy), as listed in Table 47.1. Assuming that the anode potentials were in the range of 600 to 650 mV versus SCE, and considering that the actual rates tend to be lower than the steady-state values, these predicted rates are within a reasonable range of the observed values.

REFERENCES

The corrosion rate at the anode of a macrogalvanic cell increases as its potential becomes more noble than the open-circuit potential. The shift of the anode potential in the noble direction is greater, the more noble the open-circuit potential of the cathode and the higher the cathode–anode area ratio. If the latter ratio is very large, the potential of the anode is practically the same as that of the cathode, the corrosion rate of the anode, or the penetration rate of localized corrosion being a maximum. The maximum penetration rate at the anode in a given macrogalvanic cell system can be estimated by the steadystate anodic current density at the potential at which the local anode is maintained The steady-state anodic polarization curves in deaerated 3% NaCl at 25 and 60 C are shown in Figure 47.1 [10]

Penetration Rate (mm/year)

619

mpy

0.3

12

1

40

0.9 3.5

35 140

3.5

140

10

400

1. F. Speller and F. LaQue, Corrosion, 6, 209 (1950). 2. I. Matsushima, 32nd Symposium on Corrosion and its Protection, Sapporo, Japan, Aug., 1985, p. 452. 3. I. Matsushima, Low-Alloy Corrosion Resistant Steels, Chijin Shokan Co., Tokyo, Japan, 1995, p. 155. 4. E. Usitalo, Proceedings of the 2nd International Congress on Metallic Corrosion, National Association of Corrosion Engineers, Houston, TX, 1963, p. 812. 5. E. R€as€anen and K. Relander, Scand. J. Metall., 7, 11 (1978). 6. V. Mitrovic-Scepanovic and R. Brigham, Proc. EVALMAT 89, Iron Steel Inst. Japan, 1989, p. 491; also see R. J. Brigham, M. McLean, V. S. Donepudi. S. Santyr, L. Malik, and A. Garner, Can. Met. Q., 27, 311 (1988). 7. K, Itoh, H. Mimura, T. Inoue, S. Sekiguchi, Y. Horii, and H. Kihara, J. Iron Steel Inst Jpn., 72, 1265 (1986). 8. S. Endo, M. Nagae, M. Suga, S. Wada, T. Sugino, and T. Nakano, Curt. Adv. Mater. Proc., Iron Steel Inst. Jpn., 4, 1890 (1991). 9. A. Tamada, Y. Shimizu, and I. Matsushima, Nippon Kokan Giho, No. 17, 271 (1976). 10. I. Matsushima, in Localized Corrosion, F Hine, K, Komai, and K. Yamakawa (Eds.), Elsevier Applied Science, London, 1988, p. 31. 11. I. Matsushima, paper presented at the Committee on Corrosion and Protection, Vol. 19, p. 94 (1980); reproduced in Corrosion Protection Handbook (Boshoku Gijutsu Binran), Japan Society of Corrosion Engineering, Nikkan Kogyo Sinbunsha, Tokyo, Japan, 1985, p. 181 12. M. Miyasaka, Ebara Eng. Rev., No. 137, 1 (1995). 13. I. Matsushima, Boshoku Gijutsu (Corros. Eng.), 25, 563 (1976). 14. I. Matsushima, Zairyo-to-Kankyo (Corros. Eng.), 44, 683 (1995).

48 WEATHERING STEEL T. MURATA Office of Technology Transfer, Innovation Headquarters, Japan Science and Technology Agency, Tokyo, Japan

A. Need for weathering steel B. History B1. Applications B2. Expansion of the use of weathering steels C. Alloying elements C1. Mechanistic aspects D. Monitoring rust films E. Design parameters F. Guidelines for bare weathering steel G. Advancement in weathering steel in the past decade G1. Life-cycle design specification for highway bridges G2. “Minimum-Maintenance” concept G3. Advanced weathering steel References

A. NEED FOR WEATHERING STEEL Recognition of global warming and environmental degradation has caused the direction of technology to shift toward environmentally conscious and sustainable development, with the result that increased efficiency in the use of limited resources is now a focus (e.g., recycling industrial materials). In addition, the environmental characteristics of industrial products, structures, and materials through their life cycles are being improved by life-cycle assessment (LCA) of the International Organization for Standardization (ISO) 14000s. Although LCA needs further improvement, it is a potentially useful methodology for systematically analyzing a product from the extraction of resources to eventual abandonment, including reuse, recycling, or disposal in terms of energy, in most cases,

and materials, emissions of hazardous substances, and wastes. The LCA approach is being used to clarify quantitatively the important role of steel and steel products in our society in terms of their abundant resources, economic availability in large quantities, workability, reliability, reasonable life-cycle cost, relatively low environmental burden, and recycling possibilities. Structural steel is expected to remain as one of most environmentally friendly materials available in the twentyfirst century. In addition, the recent life-cycle study of automobiles, houses, bridges and so on conducted by the Engineering Academy of Japan has demonstrated that 7085% of the energy consumption by these products throughout their life cycle occurs during service period, rather than during production or product assembly [1]. This observation indicates that materials performance, durability, and reliability are key considerations in the design of a product or a structure for its life-cycle energy efficiency. Figure 48.1 shows an example of the framework for the systematic assessment of materials performance, including (a) life-cycle cost (LCC) analysis, (b) life-cycle safety (LCS) design, (c) LCA for energy and materials, and (d) diagnosis of life prediction based on the operating mechanisms. Our living environments consist mainly of freshwater, seawater, the atmosphere, and soil. These four categories include the environments in which the majority of structural steel products and large structures are exposed. Corrosion mechanisms of structural steels and the dominant corrosion parameters in these environments are discussed in other chapters. In this chapter, carbon and low-alloy structural steels resistant to atmospheric corrosion (i.e., weathering steels) are discussed. Weathering steels have been widely accepted in recent years from the viewpoint of saving not only money but energy and resources.


621

622

WEATHERING STEEL

FIGURE 48.1. Framework for the systematic assessment of weathering steel performance.

B. HISTORY Weathering steel contains Cu, P, Cr, Ni, and Si as alloying elements totaling a maximum of a few percent. During field exposure, an adherent, compact, and protective layer of corrosion products grows under a coarse surface rust layer retarding further corrosion. Weathering steel was originally developed by U.S. Steel in 1933 as CORTEN, a high-strength, low-alloy structural steel of yield strength >35 kgf/mm2 and corrosion resistance superior to that of Cu-bearing steel [2]. The initial chemical composition is shown in Table 48.1 [3] with more recent American Society for Testing and Materials (ASTM) specifications [4]. Buck [5] reviewed investigations dating from 1908 which showed the beneficial effects of Cu as an alloying element in steels for atmospheric exposure. In 1916, ASTM took an initiative to conduct long-term exposure tests to evaluate the effect of Cu on atmospheric corrosion resistance of various steels. The final report [6] confirmed the superior corrosion resistance of steels containing Cu up to 0.3%, compared to those without Cu (37 mm (1.5 in.). Nickel was added to minimize plate damage during rolling due to enrichment of Cu in the surface scale [10] and to improve the corrosion resistance in chloride-bearing environments [2]. In the early applications, weathering steel was used for freight trains, trucks, agricultural equipment and so on with various paints. In eight-year atmospheric field exposure tests, painted weathering steel lasted 1.5–4 times longer than on conventional steels [11]. Paints, which retard corrosion initiation and minimize stains from corrosion products, are evaluated for their contribution to life-cycle economy, reduction of environmental stains, and their ease of use. Bare weathering steel was first used in a full-scale application in the John Deere and Co. office building in Moline, Illinois, in 1964 which stimulated the use of weathering steel in buildings and other structures such as Chicago Civic Center, towers for power transmission lines, and highway bridges. Thus the use of weathering steel expanded in the United States (e.g., by 1993, there were 2300 bridges in the United States without coatings) [12]. Guidelines were prepared for the removal of mill scale and contaminants, and for the design to install water drainage. B2. Expansion of the Use of Weathering Steels Weathering steel was first commercialized in the United States, and the use of weathering steel then has spread to other parts of the world where different environmental conditions exist. In the late 1950s, weathering steel was modified in Japan [3] into Cu–P, Cu–P–Ti, Cu–Cr–Ni–P (CORTEN), Cu–P–Cr–Mo and other systems. The main features included (1) diversification of steels according to P content: high P for high corrosion resistance and low P for better welding properties; (2) three grades of tensile strength covering 41, 50, and 58 kgf/mm2 (58, 71, and 83 psi, respectively) for welded structures; and (3) a variety of coatings to stabilize the corrosion products early in the exposure period to avoid stains of corrosion products [13–16]. During the 1960s, air pollution in industrial areas in Japan, up to 0.07 ppm SO2, caused by combustion of sulfur-bearing fossil fuel, along with high rainfall and humidity levels, and rather high chloride concentration near the seacoast led to poor performance of weathering steel. The alloy design was modified and coatings were developed to meet the needs for the local applications [3]. Steelmakers collaborated in 12-year exposure tests that

TABLE 48.1.

Compositions of Weathering Steelsa C

Si

Mn

P

S

Cu

Cr

Ni

Weight Percent (w/o)

CORTEN example [2] 1940 CORTEN example [2] USS CORTEN A

0.10 0.09 50.12 0.06 50.12

0.50–1.00 0.93 0.25–0.75 0.54 0.25–0.75

0.10–0.30 0.30 0.20–0.50 0.48 0.20–0.50

0.10–0.20 0.16 0.07–0.15 0.11 0.07–0.15

— 0.035 50.05 0.030 50.05

0.20–0.50 0.42 0.25–0.55 0.41 0.25–0.55

0.50–1.50 1.1 0.30–1.25 1.0 0.50–1.25

0.03 50.65 0.51 50.65

—

USS CORTEN B

50.19

0.30–0.65

0.80–1.25

50.04

50.05

0.25–0.40

0.40–0.65

50.40

V0.02–0.10

USS CORTEN B-QT

50.19

0.30–0.65

0.80–1.25

50.04

50.05

0.25–0.40

0.40–0.65

50.40

V0.02–0.10

USS CORTEN C

50.19

0.30–0.65

0.80–1.35

50.04

50.05

0.25–0.40

0.40–0.70

50.40

V0.40–0.10

ASTM A242 ASTM A588 ASTM A588 ASTM A588 ASTM A588

50.15 50.19 50.20 50.15 0.10–0.20

— 0.30–0.65 0.15–0.50 0.15–0.40 0.50–0.90

51.00 0.80–1.25 0.75–1.35 0.80–1.35 0.75–1.25

50.15 50.04 50.04 50.04 50.04

50.05 50.05 50.05 50.05 50.05

50.20 0.25–0.40 0.20–0.40 0.20–0.50 50.30

— 0.40–0.65 0.40–0.70 0.30–0.50 0.50–0.90

— 120 C) for a critical period of time. This heating also causes the alloy to become susceptible to intergranular corrosion. However, if the heating is carried out for a sufficiently extended period of time, the susceptibility to intergranular corrosion again disappears, probably because substantially all the copper has precipitated out of solid solution and therefore the zones adjacent to the grain boundaries are no more depleted in copper than are the other areas in the grain boundaries [3]. For many of the other aluminum-base alloys, metallurgical factors have relatively little effect on resistance to corrosion. Alloys such as 1100, 3003, 5052, 6053, and 6061 are relatively insensitive in this respect [3]. In many types of exposure, cold work does not appreciably affect the resistance to corrosion of a wide variety of aluminum-base alloys. In solutions of nonoxidizing acids, however, cold work stimulates corrosion to some extent and also indirectly stimulates corrosion of aluminum alloys containing over about 5% magnesium. With the latter alloys, severe cold work increases the tendency for a magnesium– aluminum constituent to precipitate from solid solution. On exposure to certain media, selective attack of this constituent then occurs [3]. The resistance to corrosion of weldments of aluminum alloys is determined, in part, by the alloy welded, by the filler alloy, and by the welding process. Galvanic cells that cause corrosion may be created because of potential differences among the parent alloy, the filler alloy, and the heat-affected zones where microstructural changes occur. Incomplete removal of fluxes after welding them may also cause corrosion [1]. Weldments in non–heat–treatable alloys generally have good resistance to corrosion. Microstructural changes in the heat-affected region in these alloys have little effect on

742

ALUMINUM AND ALUMINUM ALLOYS

potential, and the filler alloys recommended have potentials close to those of the parent alloys. In some heat-treatable alloys, however, the effect on potential of microstructural changes may be large enough to cause appreciable corrosion in more aggressive environments; the corrosion is selective, either in the weld bead or in a restricted portion of the heataffected zone. To a considerable degree, the effect of microstructural changes on corrosion in the heat-affected zone can be eliminated by postweld heat treatment. Stress corrosion cracking in weldments usually is caused by residual stresses introduced during welding, but its occurrence is rare. Brazed joints in aluminum alloys also have good resistance to corrosion. Excessive corrosion is usually caused by fluxes that are not removed completely or that are removed by a treatment that, together with the fluxes, may cause corrosion. Soldered joints have a resistance to corrosion satisfactory for applications in milder environments, but not for those in more aggressive ones [1]. D2. Composites Aluminum alloys reinforced with silicon carbide, graphite, alumina, boron, or mica show promise as metal matrix composites with increased modulus and strength and are potentially well suited to lightweight structural applications, including aerospace and military needs. The structures of continuous fiber metal matrix composites (MMCs) are equivalent to those in polymer matrix composites. Industrial applications have emerged recently, for example, reinforced pistons for assembly in light diesel engines, 12% alumina, 9% carbon fiber reinforced Al–12.7% Si MMC cylinder liner [47], and so on. Generally, long-term tests have shown that the introduction of a reinforcement phase reduces the resistance to corrosion. The extent of this reduction largely depends on the reinforcement species and, form. As with conventional aluminum alloys, the fabrication method and heat treatment infhience the corrosion resistance of MMCs and must be carefully controlled. As surface protection will be advisable in certain applications, it is encouraging to see a variety of standard techniques showing promise for MMCs. From the studies performed on the corrosion fatigue of MMCs in saline environments it appears that they are marginally inferior to their matrix alloys [46]. E. CORROSION PREVENTION AND PROTECTION E1. Design, Alloy Selection, and Joint-Sealing Compounds During conception, the corrosion specialist should identify the different types of corrosion and prevention methods. Among the most common harmful effects are galvanic

action, resulting from direct contact between aluminum and a dissimilar metal, such as copper, and indirect galvanic effects resulting from contact between aluminum and solutions containing reducible compounds of heavy metals. In some cases, design or construction will prevent serious corrosion even though no other factors are altered. Similarly, since the various alloys of aluminum differ widely in behavior, the selection of the most suitable alloy is important [3]. Corrosion can be prevented or reduced by cladding with a more corrosion-resistant alloy, such as high-purity aluminum, a low magnesium–silicon alloy, or an alloy of 1% zinc. All of these cladding materials are frequently employed to give added corrosion protection to the 2000 and 7000 series alloys. The cladding on each side is 2–5% of the total thickness. Aluminum-base alloys, such as 1100, 3300, 5052, 6053, Alclad 3300, Alclad 1017-T, and Alclad 2024-T, are highly resistant when freely exposed to most natural environments. They will all discolor or darken appreciably in most outdoor exposures but will suffer no structurally appreciable changes in properties unless exposed in relatively thin sections, that is, 0.5 mm/year, 20 mpy) regimes in these acids. The tops of the bars represent the boiling points. For perspective, the boundaries between the “very safe” and “moderately safe” regimes [i.e., the 0.1-mm/year (4-mpy) lines] are plotted in Figures 55.6 and 55.7 alongside similar lines for 316L stainless steel, 20Cb-3Ò alloy (an austenitic stainless steel designed to resist sulfuric acid), 254SMOÒ (one of the 6 wt % molybdenum-bearing stainless steels), and C-2000Ò alloy, a nickel–chromium–molybdenum material. In sulfuric acid, ULTIMET alloy is in the same performance

band as 20Cb-3 and 254SMO alloys, being far better than 316L but not as good as the Ni–Cr–Mo materials. In hydrochloric acid, ULTIMET alloy is second only to C-2000 alloy and much more resistant to this aggressive acid than the three stainless steels. The resistance to localized attack of ULTIMET alloy as compared with the Ni–Cr–Mo (C-type) alloys and austenitic stainless steels is summarized in Figure 55.8. These tests were performed according to the ASTM G 48 procedures using teflon crevice blocks. AL-6XNÒ alloy is also a 6 wt % molybdenum-bearing stainless steel. From Figure 55.8, it is evident that the resistance to localized attack of ULTIMET alloy is equivalent to that of the Ni–Cr–Mo alloys. This has been confirmed in pitting tests also. With regard to the SCC resistance of the corrosion and wear-resistant cobalt materials, such as ULTIMETalloy, they appear to be similar to the super austenitic stainless steels in their susceptibility. Reference [8], in which a four-point bend

FIGURE 55.5. Isocorrosion diagram for ULTIMET alloy in hydrochloric acid.

ALLOYS DESIGNED FOR AQUEOUS CORROSION AND WEAR RESISTANCE

FIGURE 55.6. Comparison of 0.1-mm/year (4-mpy) lines for stainless steels, a Ni–Cr–Mo alloy, and ULTIMET alloy in sulfuric acid.

FIGURE 55.7. Comparison of 0.1-mm/year (4-mpy) lines for stainless steels, a Ni–Cr–Mo alloy, and ULTIMET alloy in hydrochloric acid.

FIGURE 55.8. Critical crevice corrosion temperature (the lowest temperature at which crevice attack is experienced over a 24-h test period) in 6% ferric chloride.

755

756

COBALT ALLOYS

test results in 30 wt % magnesium chloride are presented, indicates that ULTIMET alloy is similar to 20Cb-3 alloy and that both these alloys possess much higher resistance to SCC than does 316L stainless steel.

7. P. Crook, A. I. Asphahani, and S. J. Matthews, U.S. Patent 5,002,731, 1991. 8. R. Baboian (Ed.), Corrosion Tests and Standards—Application and Interpretation, American Society for Testing and Materials, Philadelphia, PA, 1995, p. 486. 9. H. Berns and A. Fischer, Wear, 162–164, 441 (1993).

REFERENCES TRADEMARKS 1. W. L. Silence, Wear of Materials—1977, American Society of Mechanical Engineers, New York, 1977, p. 77. 2. P. Crook, Mater. Perform., 30(2), 64 (1991). 3. R. D. Gray, STELLITE—History of the Haynes Stellite Company, 1912–1972, Cabot, Boston, MA, 1974. 4. C. P. Sullivan, M. J. Donachie, Jr., and F. R. Morral, CobaltBase Superalloys—1970, Centre d’Information Du Cobalt, Brussels, Belgium, 1970. 5. C. T. Sims and W. C. Hagel, The Superalloys, Wiley, New York, 1972. 6. Metals Handbook, 9th ed., Vol. 9, Metallography and Microstructures, ASM, Metals Park, OH, 1985.

STELLITEÒ and TRIBALOYÒ are registered trademarks of Deloro Stellite. HAYNESÒ , ULTIMETÒ, HASTELLOYÒ, and C-2000Ò are registered trademarks of Haynes International, Inc. VITALLIUMÒ is a registered trademark of Howmedica. NITRONICÒ is a registered trademark of Armco Steel. 20Cb-3Ò is a registered trademark of Carpenter. 254SMOÒ is a registered trademark of Avesta Sheffield. AL-6XNÒ is a registered trademark of Allegheny Teledyne, Inc. MAR-MÒ is a registered trademark of Lockheed Martin.

56 COPPER AND COPPER ALLOYS C. A. C. SEQUEIRA Instituto Superior Tecnico, Lisboa, Portugal

A. Introduction B. Pitting corrosion B1. Pitting mechanisms B2. Pitting prevention C. Stress corrosion cracking C1. Mechanisms C2. Effects of alloy composition C2.1. Aluminum bronzes C2.2. Copper nickels C2.3. Silicon bronzes C2.4. Tin brasses C2.5. Tin bronzes C2.6. Nickel silvers C2.7. Aluminum brasses C2.8. Manganese bronzes C2.9. Brasses C2.10.Copper–gold alloys C3. Specific corrodents C3.1. Ammonia system C3.2. Sulfur dioxide C3.3. Citrate and tartrate solutio C3.4. Pyridine and ethylenediamine solution C3.5. Other environments D. Atmospheric corrosion E. Corrosion in waters E1. Copper E2. Copper alloys E2.1. Brasses E2.2. Copper–tin alloy E2.3. Aluminum bronze E2.4. Silicon bronzes E2.5. Copper–nickel alloys E3. Microbial corrosion E4. Biofouling E5. Condensers and heat exchangers F. Corrosion in gases F1. Oxidation of copper

F2. Oxidation of copper alloys F3. Corrosion by gases other than oxygen References

A. INTRODUCTION Copper is the dominant material for domestic water systems in Europe, North America, Australasia, and many countries of the Commonwealth. The developing world is now rapidly expanding its use of copper tubing as their societies demand increasing quantities of disease-free water. The annual world production of copper water tubing is 500,000 tonnes, equivalent to 1.25 billion m or 0.75 million miles. The reasons for this enormous consumption of copper plumbing tube are its excellent corrosion resistance, its ease of fabrication during installation, and hence lowinstallation costs, and additionally, its contribution to health and the maintenance of healthy water. Further, copper is seen as environmentally friendly due to its potential to be 100% recycled. Failures of copper water tube from corrosion are rare but in general well-understood events. The high level of quality control exercised in modern tube producing plants plus copper’s excellent corrosion resistance gives failure rates of 99% Cu 99% Cu > 94% Cu Cu–Zn–Sn–Pb (75–89% Cu) Cu–Zn–Sn–Pb (57–74% Cu) Cu–Zn–Mn–Fe–Pb Cu–Zn–Si Cu–Sn–Zn–Pb Cu–Ni–Sn–Zn–Pb Cu–Al–Fe–Ni Cu–Ni–Fe Cu–Ni–Zn–Pb–Sn Cu–Pb

Cu–Ni–Fe Cu–Ni–Zn

Note: UNS ¼ Unified Numbering System.

Pure copper is a very soft, malleable metal. It is alloyed with small quantities of metals such as Be, Te, Ag, Cd, As, and Cr to modify the properties for particular applications, while retaining many of the characteristics of the pure metal. Much larger alloying additions of Zn, Sn, and Ni are made to improve the mechanical properties of the metal, and to retain its excellent corrosion resistance under more arduous service conditions. Nickel permits increased flow rates in water systems; zinc gives increased resistance to sulfide attack. Typical groups of copper alloys that find application in many environments are listed in Table 56.1. Apart from the wide use of copper in freshwater supply lines and plumbing fittings, copper and its alloys provide superior service in many other applications. They are used for structures open to the atmosphere, for example, in architecture and sculpture. They are used immersed in freshwater and seawater heat exchangers and condensers, as well as in industrial, chemical, and power-generating plants, and buried in the earth for water distribution systems. In the open air, copper forms a green patina that, in its most stable form, consists of basic copper sulfate, CuSO43Cu(OH)2, although in marine environments it

may contain chloride, or carbonate in industrial areas. This decorative long-lasting coating makes copper an ideal material for low- maintenance roof coverings and for gutters and channels. A small amount of copper dissolves in water that runs over the metal surface, and this can precipitate on to other less noble metals downstream in the water cycle, leading to galvanic corrosion. Cast iron gutters and pipes used in conjunction with copper roofs benefit from a bituminous or other impervious coating to reduce the possibility of galvanic corrosion. General corrosion or stress corrosion cracking (SCC) may become a problem in industrial areas if ammonium compounds are present in the atmosphere. Copper and its alloys can safely be buried in most soils, although high corrosion rates have been experienced in those containing cinders or acid peat. If it is expected that corrosion will be a life-limiting factor in the use of the material, it can be protected with bituminous, plastic, or paint coatings. Dezincification can be a problem in brasses with high zinc levels, and it is best to avoid the use of these alloys unless they are specifically required to counter the difficulties that may result from high sulfide levels in the soil.

INTRODUCTION

Copper and its alloys are used extensively in seawater distribution systems and in treatment units, condensers, and heat exchangers where fresh or salt water is used for cooling. Many components in valves, pumps, and taps, as well as pipes and pipe fittings, are made from copper alloys. A distinction is often made between corrosion in freshwater and in seawater. However, the same types of corrosion problem are found in both environments, and such a clear-cut distinction cannot be made. Rather, the change from pure to salt water with varying degrees of pollution should be regarded as a gradually increasing aggressiveness in the environment, aggravated by increases in flow rates and changes in the temperature and oxygen content of the water. The main problems of copper alloys in water systems are differential aeration corrosion, erosion corrosion, SCC, and demetallification. Differential aeration corrosion is mainly a design problem, although pitting may occur under very slow flow rates, which starve the metal surface of oxygen. Erosion corrosion is a function of flow rate. Pressure changes in a liquid on passing through valves and pumps give rise to cavitation damage, while entrained air or abrasive particles disrupt protective surface films to produce shallow, horseshoe-shaped pits. The deterioration can be very rapid. Ammonia and its salts, together with mercury-based compounds, are the prime cause of SCC. Dealloying affects many of the alloys, the commonest being dezincification of brasses containing > 15% zinc, although dealuminification and denickelification have been reported for aluminum bronzes and cupronickels. Dealuminification is most prevalent in aluminum bronzes containing the g-2 phase in the microstracture, and is most serious when the g-2 forms a continuous grain boundary network. Rapid cooling from > 600 C (1115 F), additions of 1–2% iron, or more than 4.5% nickel, control dealuminification, but microstructural changes that occur during welding can still lead to corrosion problems in the heat-affected zone (HAZ) around the weld. Reheat treatment to remove unsatisfactory postweld microstructures can pose serious problems to the engineer; heating, handling, and quenching large units without introducing distortion while achieving uniform properties throughout the material may be impossible. Brass fittings may dezincify, especially when the b phase is present. The loss of zinc is accelerated by hightemperature, increased chloride content, low flow rates, and differential aeration. Additions of 1% tin and 0.04% arsenic, phosphorus, or antimony inhibit dezincification. However, phosphorus can lead to intergranular corrosion, and most manufacturers use arsenic as an inhibitor in brasses. Inhibited a brasses are immune to dezincification in most waters, but the effect of tin and arsenic additions to a/b brasses is not predictable in controlling dezincification. There have been many cases of dezincification in the duplex brasses in both fresh and seawater. In some instances in

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potable water distribution systems, duplex fittings that have given many years’ service suddenly begin to lose zinc when only a slight change in water chemistry occurs. A protective film of carbonate may be deposited on the metal surface from water containing carbon dioxide and oxygen. As the flow rate increases, copper and brass tubes become more prone to impingement attack. Aluminum brass and cupronickel offer a greater resistance to higher flow rates, but both have maximum limits that must not be exceeded or the surface film on the metal will be destroyed. While the maximum velocity for inhibited Admiralty brass and aluminum brass is lower than that for cupronickels, they both give better service should sulfide be present in the water, either as a pollutant in rivers or estuaries or in chemical and oil production plants. However, work is in hand to develop coatings for copper-based alloy condenser tubes in landbased and marine systems. These coatings will offer protection against sulfide-polluted water while the natural film is established. Experimental programs have shown that these coatings will also be effective on both cupronickel and aluminum brass tubes [1]. Tin bronzes and phosphor bronzes have good resistance to flowing seawater. The alloys containing 8–12% tin are less susceptible than brasses to SCC and have excellent resistance to impingement attack and to attack in acid waters. The aluminum oxide film on both aluminum brass and aluminum bronze, which confers the additional corrosion protection, reduces the dissolution rate of copper ions from the alloy and makes them less effective in resisting biofouling. The resistance of some copper alloys to erosion–corrosion is improved when small quantities of iron are present in the alloy or the water. The iron apparently produces a tougher surface film. This has led to the use of iron sacrificial pieces, in preference to the normal zinc sacrificial anodes, in the water boxes of condensers and heat exchangers that use copper-based tubes and tube plates. The iron ions are absorbed into the film from the water and confer the additional resistance to impingement attack. Zinc ions do not have this beneficial effect. A very clear distinction can be made between acids that can be safely handled in copper-based equipment and those that cause catastrophic attack. Nonoxidizing acids, such as acetic, phosphoric, dilute sulfuric, and hydrochloric, can be safely handled providing the concentrations of oxidizing agents such as entrained or dissolved air, chromates, and iron(III) ions are kept very low. Oxidizing acids, nitric or concentrated sulfuric, and those containing oxidizing agents must not be handled in copper-based systems. Small additions of oxidizing agents are particularly dangerous in hydrochloric acid, causing a dramatic increase in the rate of metal loss. Before using copper alloys in acid systems, tests should always be made with the particular liquid to be

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COPPER AND COPPER ALLOYS

processed, reproducing the actual plant conditions as closely as possible. In general, copper and its alloys are resistant to attack by alkalis except ammonium hydroxide and those containing ammonium or cyanide ions. Ammonium ions promote SCC and both ions form complex species such as [Cu(NH3)4]2 þ and [Cu(CN)4]2 , which do not allow the double layer to develop to polarize the corrosion cell, hence the corrosion rate remains high. Iron(III) and tin(IV) salts are aggressive to copper alloys. Copper itself suffers general corrosion, to thin the cross section, in sulfur compounds.

B. PITTING CORROSION B1. Pitting Mechanisms Since the first description of pitting of copper tubes in contact with water in 1950 [2], extensive investigations/have since been conducted with various natural waters, reflecting the practical importance of this problem [3–12]. Not only copper, but also copper alloys, such as brasses [13, 14], bronzes [15], some cupronickels (e.g., 70 Cu–30 Ni [16–18]), and other alloys [19–22] can be damaged by pitting. Pitting of copper does not occur exclusively in chloride-containing solutions but in hydrogen carbonate solutions as well [23– 25]. Kinetics of passivation and pitting corrosion of copper have been extensively studied using electrochemical techniques [26–30]. These techniques include measurement of redox potentials and of the potentials of corroding electrodes [31], measurement of current density between coupled electrodes, polarization resistance, potentiokinetic curves [32], impedance measurements, and electrochemical noise [30, 33]. Other methods of investigation include eddy current examinations [34], morphological studies [35], and use of chemical and biochemical microsensors [36]. Anodic polarization curves of copper in NaCl solutions show that the breakdown potential is lower the higher the NaCl concentration. When copper specimens are immersed for long periods of time in dilute NaCl solutions, pitting is not observed, and general corrosion occurs. The presence of other halogen anions in NaCl inhibits pitting of copper, but the action of I differs from that of Br [37]. According to Mor and Beccaria [37], with iodides inhibition is caused by the formation of a Cul layer on the metal surface, whereas KBr acts by reducing the oxygen content of the solution. In very dilute KBr solutions, however, copper corrosion increases. In hydrogen carbonate solutions, with increasing NaHCO2 concentration, the resistance of the oxide layer toward general corrosion decreases [25], which is probaby caused by increasing soluble copper carbonate complex

formation. However, increasing HCO3 oncentration promotes break down of the oxide layer. As the temperature is increased from 25 to 90 C (80 to 197 F), the oxide film is less protective and the pitting potential shifts to less positive values. Kristiansen [38] studied pitting of copper in distilled water containing 10 mg/L SO4 2 plus 5 mg/L CO2, with and without Fe3 þ additions, at temperatures of 45, 50, and 60 C (116, 125, and 143 F). In aerated neutral water ferric ions are reduced on copper surfaces according to Fe3 þ þ Cu ¼ Cu þ þ Fe2 þ Ferrous ions are reoxidized to Fe3 þ , causing further copper corrosion. Using radioactive 59 Fe, it was established that pits nucleated where iron was present on the metal surface. It was also found that the highest rate of corrosion occurred at 50 C (125 F), which was explained by the decomposition of the basic copper carbonate deposit and formation of a more protective copper oxide at higher temperatures. Pits were also found on copper when no iron had been added, although the presence of iron in the water could not be excluded. The deleterious effect of ferric ions has also been reported by Molvina et al. [39]. The problem of copper tube pitting in supply waters still persists in many countries. This type of localized corrosion was called “nodular pitting” by Campbell [15] because corrosion produces small areas of deep attack covered by small mounds or nodules of corrosion products. According to Mattsson [40], three types of pitting can be distinguished for copper: Type 1 occurs on annealed or halfhard tubes in cold tap water, caused by a continuous carbon film formed on the inner tube surface during bright annealing; type 2 occurs on hard drawn tubes in hot tap water of low (60 C (143 F) [15]. Mattsson [40] found that pitting occurs in low pH water (generally pH 5–7) and that water that causes pitting has a relatively low HCO3 content (100 mg/L), whereas no pitting occurs in water with a higher HCO3 content (100–300 mg/L). The SO42 concentration is 15–40 mg/L. At higher HCO3 concentrations, the protective basic copper carbonate forms. Mattsson [40] observed that in carbonate containing water (15–70 mg/L HCO3 plus >0.2 mol/L SO42 , a basic copper sulfate crust forms that covers the capillary mouth and thus creates an occluded cell at the anode where the reaction can occur. 2Cu þ H2 O ¼ Cu2 O þ 2H þ þ 2e

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As in type 1 pitting, pits initiate below the metal surface deposit. Mattsson also found that in tubes with pitting, the inner tube wall generally had a greenish gray surface coating, often with a high aluminum content and containing basic copper carbonate; a green, basic copper sulfate crust covered the pits. Cornwell et al. [44] suggested the following mechanism of copper pitting in aerated supply waters: Cu þ þ Cl ¼ CuCl CuCl hydrolyses to form Cu2O, which is precipitated on the metal surface: 2CuCl þ H2 O ¼ Cu2 O þ 2HCl The cathodic reaction supporting the anodic dissolution process is oxygen reduction: O2 þ 2H2 O þ 4e ¼ 4OH For corrosion to proceed, the hydroxyl ions produced at the cathodic sites must be removed. This occurs more rapidly in acid supply water or water that contains bicarbonate ions: OH þ HCO3 ¼ CO3 2 þ H2 O The final reaction causes precipitation of mixed calcium carbonate and basic copper carbonate scale. B2. Pitting Prevention There are several ways of avoiding pitting corrosion of copper in waters. Ferrous ion injection has been used for many years in Europe and Japan to protect aluminum, brass, and cupronickel tubes, mainly in polluted waters [47–50]. The beneficial effect of ferrous ion, as FeSO4, can be explained on the basis of either of the following mechanisms: 1. Electrochemical protection: Following Cornwell et al. [41], Lecointre et al. [51] and Pourbaix [52] observed that if the tube surface can be maintained at an electrode potential lower than the protection potential of copper, existing pits cannot grow and new pits cannot form. The redox couple Fe2 þ –FeOOH can maintain the tube surface below this protection potential, and this protection process should remain valid for copper alloys (e.g., brass) and waters containing chloride (e.g., seawater and brackish water). Based on the protection potential for copper alloys in chloride-containing waters and on the electrochemical action of the addition of ferrous ions, a method of control is proposed: The potential of the tubes is measured and ferrous ions are

injected to maintain this potential below the protection potential. 2. Formation of a protective film: Following North and Pryor [53], ferrous ions are oxidized. directly in the water by dissolved oxygen to form colloids. By a process of electrophoresis, these colloids adhere to the walls of the tubes to form a layer, consisting primarily of FeOOH lepidocrocite, which is considered to be protective. According to Gasparini et al. [54], ferrous ions precipitate on local alkalized cathodic areas to form hydroxides that are subsequently oxidized to produce a layer of lepidocrocite. North and Pryor [53] considered that the film formed on copper in a NaCl solution containing a small amount of FeSO4 is mainly composed of h-FeOOH; Cu2O was also detected under the film. Castle et al. [55] and Epler and Castle [56] used ESCA to characterize the protective films formed on copper-based condenser tubes protected by ferrous sulfate injection. The film, was found to be composed mainly of two superimposed layers: an inner layer, white, containing hydrotalcite, Mg6Al2(OH)6CO3 4H2O, or other compounds of the hydrotalcite family and paratacamite, Cu2(OH)3Cl; and an outer layer, brown, porous, which is, a mixture of lepidocrocite (essentially h-FeOOH), paratacamite, iron oxides, and copper oxides. Hydrotalcite can exercise a buffer effect through the effect of Mg2 þ ions. In waters containing phosphates, tricalcium phosphate is found in the film. In polluted waters containing sulfides, Cu2S is found in the external layer of the film. Assuming that the protection is given by a film, one method of monitoring the thickness, porosity, arid other characteristics of the film is to perform polarization tests [52, 53]. Ultraviolet (UV) photochemical decomposition of residual chlorine [57], adequte metal and water treatment (e.g., inhibitors for the solution or alloying additions to the metal tube) [58–63], the application of artificial protective films [64], and control of the redox potential have all been recognized as means of suppressing or preventing pitting corrosion of copper alloy tubes in waters. But, of course, an important effective means of combatting or avoiding pitting attack is by appropriate alloy selection. Of the copper alloys, the most pit resistant are the copper–aluminum alloys known as aluminum bronzes with 5% Sn are especially resistant to impingement attack. In general, the tin bronzes are noted for their high strength. Their main application is in water service for such items as valves, valve components, pump casings, and so on. Because of their corrosion resistance in stagnant waters, they also find wide application in fire protection systems. Tables 56.5 and 56.6 list the principal tin bronzes used for corrosion engineering. C2.6. Nickel Silvers. Nickel silvers are Cu–Zn–Ni alloys. Those containing 5, 10, and 15% Ni did not crack in mercurous nitrate solution after a week, but they can crack if stresses are sufficiently high, usually well above the yield point [122]. The most common nickel silvers are C75200 (65 Cu–17 Zn–18 Ni) and C77000 (55 Cu–27 Zn–18 Ni). They have good resistance to corrosion in both fresh and saltwaters. Primarily because their relatively high nickel contents inhibit dezincification, C75200 and C77000 are usually much more resistant to corrosion in saline solutions than brasses of similar copper content.

C2.2. Copper Nickels. Copper–nickel alloys are not very susceptible, 70 Cu–30 Ni cold-worked tubes do crack on exposure to mercurous nitrate after extremely heavy hollow sinking. Because of its high resistance to SCC, 70 Cu–30 Ni alloys have replaced alloys susceptible to SCC, such as Admiralty brass in environments containing small amounts of NH3. It has been found that susceptibility to SCC of Cu–Ni–Si alloy in NH3 atmosphere depends on aging time and on the extent of plastic deformation preceding or following aging, marked improvement in resistance being caused by prior cold working and to a lesser extent by deformation subsequent to aging. The beneficial effect of plastic deformation is attributed to the increase in the number of sites for preferential corrosion and consequently decreasing the rate of penetration [119, 120].

C2.7. Aluminum Brasses. Precipitation of CuAl2 at grain boundaries and formation of Cu-depleted zones adjacent to the boundaries causes a potential difference of 200 mV in NaCl solution between the boundary and the body of the

TABLE 56.5. Some Common Wrought Copper–Tin Alloys: Chemical Composition UNS Number C-51000 C-51100 C-52100 C-52400 C-54400

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Cu

Pb

Fe

Sn

Zn

P

Rem Rem Rem Rem Rem

0.05 0.05 0.05 0.05 3.5–4.5

0.10 0.10 0.10 0.10 0.10

4.2–5.8 3.5–4.9 7.0–4.9 9.0–11.0 3.5–4.5

0.30 0.30 0.20 0.20 1.5–4.5

0.03–0.35 0.03–0.35 0.03–0.35 0.03–0.35 0.01–0.50

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TABLE 56.6. Some Common Cast Copper–Tin Alloys: Chemical Composition UNS Numbers C-90300 C-90500 C-92200 C-93700 C-93800 C-93900 C-94700

Cu

Sn

Pb

Zn

Fe

Sb

Ni

S

P

86.0–89.0 86.0–89.0 86.0–89.0 78.0–92.0 75.0–79.0 76.5–79.5 85.0–90.0

7.5–9.0 9.0–11.0 5.5–6.5 9.0–11.0 6.3–7.5 5.0–7.0 4.5–6.0

0.30 0.30 1.0–2.0 0.8–11.0 13.0–16.0 14.0–18.0 0.10

2.0–5.0 1.0–3.0 2.0–5.0 0.8 0.8 1.5 1.0–2.5

0.20 0.25 0.25 0.15 0.15 0.4 0.25

0.20 0.20 0.25 0.55 0.8 0.50 0.15

1.0 1.0 1.0 1.0 1.0 0.8 4.5–6.5

0.05 0.05 0.05 0.08 0.08 0.08 0.05

0.05 0.05 0.05 0.15 0.05 1.5 0.05

grain, and this potential difference is one cause of intergranular cracking of aluminum brasses [123–126]. These alloys are immune to SCC when quenched from the solution heat treatment temperature. Heating the alloy to 190 C resulted in precipitation, at the grain boundaries of CuAl2, which is cathodic to the grain, and formation of a precipitate-depleted zone adjacent to the grain boundaries, thus developing an electrochemical cell [127–130].

and Tracy [108] were able to crack brasses containing only 5.18% Zn and Logan [136] easily cracked very large grained brass containing 10% Zn. Elements such as lead, tellurium, beryllium, chromium, phosphorus, and manganese have little or no effect on the corrosion resistance of copper–zinc alloys. These elements are added to enhance such mechanical properties as machinability, strength, and hardness. Table 56.7 lists the compositions of some of the brasses that are used in corrosion engineering. Brasses C44300–C44500, known as Admiralty brasses, are resistant to dealloying because of the tin in the alloy. Admiralty brass is used mainly in the handling of seawater and freshwater, particularly in condensers. Because these brasses are resistant to hydrogen sulfide, they are used in petroleum refineries. Red brass, an alloy containing 15% zinc, has basically the same corrosion resistance as copper, but with greater mechanical strength. Most brass piping and fittings are produced from this alloy.

C2.8. Manganese Bronzes. Manganese bronzes are susceptible to SCC in NH3 atmospheres, but less susceptible than Cu–Zn alloys [131]. The time to cracking depends on composition, the minimum in the cracking time versus composition curve lying near 5–6% Mn. The mode of cracking was also altered with changing composition. Alloys containing 4.94% Mn failed intergranularly, whereas those containing 21.38 and 24.5% Mn failed transgranularly. Alloys containing 10.64% Mn showed mixed mode of cracking. With increasing Mn content, alloys became considerably more resistant in seawater and NaOH, but the resistance fell sharply in H2S04 when Mn content exceeded 30% [107]. Alloys containing 11.8 and 24.2% Mn corroded in H2S04 at the same rate as Cu.

C2.10. Copper–Gold Alloys. Copper–gold alloys belong to a novel but important alloy system with specific sets of properties. The Cu–Au alloys show an anodic polarization behavior that is typical of binary alloy systems with a relatively large difference in the standard potentials of the two alloy components. The current response is the result of copper electrodissolution. Another major advantage of the Cu–Au alloy system for electrddissolution and corrosion studies is the fact that hydrogen cannot be formed on the

C2.9. Brasses. Brasses are extremely susceptible to SCC and have been widely investigated in various environments [132–135]. Brasses containing 70% Cu cracks were predominantly intergranular. Three main environmental factors contribute to the specificity of NH3 in causing SCC of copper alloys: 1. Corrosion product films reduce the overall anodic activity of the surface while maintaining good electron transport properties to facilitate cathodic process. 2. Enhanced preferential anodic dissolution of Zn at grain boundaries, associated with tarnishing and crystallographic features, leads to “chemical differentiation” of the surface. 3. Stabilization of the Cu(I) valency state by complex formation introduces a very effective cathodic reaction system. Maximum susceptibility to cracking occurs when all these factors operate together [92, 151]. C3.2. Sulfur Dioxide. In early research, SO2/water/air cracked brass [145, 152]. In damp air containing SO2, SCC of brass occurs if the concentration of SO2 is between 0.05 and 0.1%, whereas at higher concentration, general overall

FIGURE 56.4. The E–pH diagrams for the Cu–BTAH–H2O systems: (a) no BTAH, (b) in the presence of 10 4 total activity of dissolved BTAH species, and (c) in the presence of 10 2 total activity of dissolved BTAH species [154]. (Reproduced by permission of The Electrochemical Socity, Inc.)

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and (c) SO2 is >10 times more efficient than O2 as a cathodic depolarizer. Tromans and Nutting [155] found that cracks in 70 Cu–30 Zn exposed to moist SO2 were intergranular and grew by linkage of microcracks at grain boundary pits. C3.3. Citrate and Tartrate Solution. The SCC of brass was found to occur in citrate solution at pH 10.3 after 84 h of immersion in 0.7 M copper citrate and 0.6 M potassium citrate [156]. In tartrate solution intergranular cracking occurred at pH 13 after 31 days immersion in solution containing 0.04 M Cu and 0.5 M potassium tartrate. Cracking is associated with Cu–citrate and Cu–tartrate complexes, the presence of which is controlled by the pH of the solution, concentration of citrate and tartrate ions, and the redox potential of the solution. Alloy C72000 is sensitive to intergranular SCC in citrate solutions containing dissolved copper in the pH range of 7–11 [157]. C3.4. Pyridine and Ethylenediamine Solution. In the presence of moist air, amines can cause intergranular SCC of stress brass, and primary amines are more damaging than secondary or tertiary amines [158]. The corrosion rate of Al–brass exposed to steam condensate containing high concentration of amines is nearly five times greater than that observed in condensate containing no amines [159]. Trace amounts of NH3 produced by the degradation of the amines might be responsible for the observed high rate of anodic dissolution in the presence of O2. Alloys C26000 and C68700 are susceptible to SCC in solution of amines [160, 161].

FIGURE 56.5. The E–pH diagrams for the Cu–BTAH–Cl–H2O systems containing 0.67 activity of [Cl 1] (equivalent to 1 M NaCl): (a) no BTAH, (b) in the presence of 10 4 total activity of dissolved BTAH species, and (c) in the presence of’10 2 total activity of dissolved BTAH species [154]. (Reproduced by permission of The Electrochemical Socity, Inc.)

effects of potential, pH, Cl , and BTAH concentrations (activities) on corrosion inhibition. It has been demonstrated by polarization studies of Cu, Fe, and Zn carrying thin condensed moisture films in the presence of O2 and SO2 that (a) Cu passivates readily, while Zn remains active, (b) the efficiency of cathodic depolarization increases with decreasing film thickness,

C3.5. Other Environments. Monel (Ni–Cu alloy) is susceptible to SCC in most aerated vapor containing HF. Alloys with 15 and 30% Cu suffered severe SCC with both intergranular and transgranular cracks [107]. The susceptibility to cracking varies with the Ni/Cu ratio. Alloys with 66% Cu resisted cracking in 14-day tests, the only local attack being shallow intergranular corrosion, but not cracking. The insolubility of nickel fluoride keeps the nickel ion concentration low and makes Ni anodic to Cu. Metallic Cu was observed on the specimen surface and also deep within the cracks. Holberg and Prange [162] reported SCC in yellow brass at 64 C (150 F) in HF alkylation in a petroleum refinery. Failure of a 70 : 30 brass chain occurred in HNO3 vapor, possibly because of the breakdown of HNO3 to NH3 [163]. The SCC of brass in an air conditioning unit was attributed to ammonium ion formed from the oxides of nitrogen produced by the corona discharge in the high-voltage precipitator unit [164]. Neither a pure solution of NaCl nor seawater cause SCC of a brass, while b brass cracks intergranularly in these media. Alloys C26000 [165] and C44300 [166] have lower fracture stresses in NaCl solutions when the metal is anodically polarized.

ATMOSPHERIC CORROSION

Failure of Cu–Au alloys in the presence of either aqueous ferric chloride or aqua regia is attributed to the formation and rapture of a mechanically weak surface film consisting of a gold-rich phase that is formed by the preferential dissolution of the less noble elements [107]. Effect of other environments (e.g., cupric acetate, polluted atmospheres, sodium chlorate solutions, sodium formate solutions, sodium hydroxide solutions, sodium nitrate solutions, sodium nitrite solutions, solder, sodium sulfate solutions, sulfide solutions, and sodium tungstate solutions) on SCC of alloys C26000 [167–169], C44300 [166], C70600 [169], and brass [168], Admiralty brass and other copper alloys are well documented in the literature [170–179].

D. ATMOSPHERIC CORROSION In addition to oxygen, moisture, and gaseous pollutants, such as sulfur dioxide (SO2), nitrogen oxides (NOx), chlorine gases (HCl, Cl2), ammonia (NH3), and ozone, airborne ionic dust particles strongly affect conditions on corroding surfaces. During atmospheric corrosion of copper, a patina is formed over time. Initially, a layer of Cu2O, CuO, and CuOxH2O forms in which cuprite (Cu2O) is the main component. Later, a patina with several corrosion products is formed. The basic copper sulfates, posnjakite, Cu4SO4(OH)6H2O, brochantite, Cu4SO4(OH)6, and antlerite, Cu3SO4(OH)4, are primary constituents. Basic copper chloride and carbonate are also found [180–184]. The mechanism of basic copper sulfate formation has been the subject of several studies that have been reviewed [183, 184]. In metropolitan areas, most of the sulfur acquired by surfaces is not in gaseous form by the reaction with SO2 but as dry deposition [185, 186]. Outdoor exposures of copper show that the main factors influencing the weight gain rates are relative humidity and concentration of aerosol particles [187]. The most abundant ions found in fine particles are SO4 2 and NH4 þ , with the ratio typically being between that of NH4HSO4 and (NH4)2SO4 [188]. The effect of submicron (NH4)2SO4 particles on the corrosion of copper at varying relative humidities (RHs) and-temperatures was studied by Lobnig et al. [180, 181], who showed that (NH4)2SO4 particles lead to the corrosion products found in natural patinas. In a recent study [189], it was investigated whether NH4HSO4 and (NH4)3H(SO4)2 may be responsible for formation of basic copper sulfate in the naturally formed patina on copper during atmospheric corrosion. At room temperature, only mixed ammonium copper sulfates, (NH4)2Cu(SO4) (OH)2xH2O, and cuprite were found up to 28 days of reaction time. It was concluded that acid ammonium sulfate is less likely than ammonium sulfate, the sulfur containing species leading to formation of basic copper sulfates in copper patinas.

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Basic copper sulfates are the dominant phases in most copper patinas, but particularly at the seacoast, the copper patina also consists of Cl-containing species. The principal source for chlorine in the atmosphere is sea salt. In marine areas, aerosols are formed with an initial chloride concentration of seawater (0.4 M), but accumulation on surfaces may result in higher concentrations. Sea salt aerosols may be scavenged by water, occurring in rain, snow,. fog, and dew, and may be far removed from the coast [182]. Chloride concentrations in precipitation are reported in the range 0.04–4 mM [185]. Another source of chloride is the use of NaCl for deicing, implying that large amounts of chloride may be found on objects close to roads. Hydrogen chloride is the dominant inorganic gaseous chlorine compound in the atmosphere [185]. The main source is reported to be the reaction of acidic trace gases with sea salt aerosols, for example NO2(g) and NaCl(s) forming HNO3 and HCl(g) [190, 191]. In addition, HC1 is emitted from the combustion of fossil fuel and refuse [185]. Concentration of HCl(g) in the atmosphere range from very low values up to 4 ppb [185]. The metal loss of copper coupons in a marine atmosphere has been reported to be 600–700 mg/cm2-year, averaged over an 8-year exposure period, the initial corrosion rate being considerably higher [192]. Copper turns brownish during the first year of exposure, while signs of green patina may be observed after 6–7 years in this environment [193]. After decades of exposure, a patchy and streaked appearance with alternating black and light green areas tends to appear. Rainsheltered areas tend to be black while areas washed by rain tend to be light green. The polymorphs [Cu2(OH)3Cl], atacamite and paratacamite in combination with different hydroxy sulfates and cuprite (Cu2O) are dominant in the black sheltered areas, while rain-washed green areas mainly contain brochantite [Cu2(OH)6SO4] and cuprite [194]. The patchy black and green appearance is reported to have become especially evident on outdoor objects since the 1960s [195]. Chlorides have been identified in copper corrosion products in marine [196] and inland [197] locations, as well as indoors [198]. Atacamite has been reported to occur more frequently than paratacamite in outdoor environments [194, 196]. Thermodynamic calculations predict atacamite to be stable at the chloride concentrations present in seawater, while it is only marginally stable at concentrations typical of rain and fog (0.4–4 mM) [195]. Jambor et al. [199] recently reported on clinoatacamite, a new mineral polymorph of Cu2(OH)3Cl. The published Joint Committee on Powder Diffraction Standards (JCPDS) Powder Diffraction File for paratacamite (PDF 25-1427) [194] is only slightly different from the diffraction pattern for clinoatacamite and is claimed to pertain to the latter compound. Thus, previous reports on paratacamite should probably be assigned to clinoatacamite instead.

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A few laboratory studies deal with the effect of chloride on the atmospheric corrosion of copper. Working with very high pollutant concentrations (percent levels), Vernon [171] concluded that HCl was more deleterious than SO2 toward copper at 50% relative humidity (RH). Feitknecht [200] proposed mechanisms for the interaction of chlorides with copper and some other metals. Eriksson et al. [201] reported that the addition of small amounts of sodium chloride caused a marked increase in the corrosion rate of bronze in humid air. More recently, Strandberg and Johansson [202] investigated the effect of NaCl in combination with O3 and SO2 on the atmospheric corrosion of copper. Large amounts of cuprite (Cu2O) formed in all environments at 70 and 90% relative humidity. The corrosive effect of salt was strong in pure humid air and in air containing O3 or SO2. Corrosion rate was correlated to the amount of chloride applied to the surface and to humidity. In an atmosphere containing a combination of SO2 and O3 at 90% RH, corrosion was rapid in the absence of NaCl. In this case, small additions of NaCl resulted in a marked decrease in corrosion rate. In the absence of SO2, tenorite (CuO), nantokite (CuCl), clinoatacamite [Cu2(OH)3Cl], and malachite [Cu2(OH)2CO3] were identified. In the presence of SO2, brochantite [Cu4(OH)6SO4], soluble sulfate, and an unknown phase occurred, while no tenorite or malachite was formed. The combination of SO2 and O3 resulted in the formation of antlerite [Cu3(OH)4SO4] and Cu2.5(OH)3SO42H2O as well. Although the severity of the atmospheric corrosion of copper alloys depends on the atmospheric contaminants, the corrosion rate usually decreases with time. In general, copper alloys are very suitable for atmospheric exposure. Highcopper alloys, silicon bronze, and tin bronze corrode at a moderate rate, and brass, aluminum bronze, nickel silver, and copper nickel corrode at a slower rate. The copper alloys, most widely used in atmospheric exposure are C11000, C22000, C38500, and C75200. Additional information on the atmospheric corrosion of copper and its alloys is available in [203–212].

E. CORROSION IN WATERS E1. Copper The corrosion resistance of copper is due to its being a relatively noble metal. Its satisfactory service in waters depends on the formation of relatively thin adherent films of corrosion products (e.g., cuprous oxide and basic copper carbonate). It has only a weak tendency to passivation, and hence the effect of differential aeration is very slight. However, the influence of copper ion concentration on the potential of copper in solution is very marked. For this reason, when there are varying solution velocities over a copper

surface (e.g., when the solution is stirred), the parts exposed to solution with the higher rate of movement become anodes and not cathodes as would be the case with iron, for example. Cuprous oxide (Cu2O) is the corrosion product predominantly responsible for protection. This oxide possess semiconducting properties and, in high-temperature corrosion studies (oxidation in the thick-film range) is usually described as a p-type semiconductor in which defects are uniformly distributed in the oxide. In waters, cuprous oxide is produced by anodic oxidation, usually under dynamical flow. That is, it is formed on the metal side and dissolved on the solution side. Under these conditions, the oxide structure is not homogeneous [213–215]. The anodic oxide thickness is very small, typically 500 nm, in the copper–water system. The coppers–water interface is really a complex copper–oxide–water system, involving solid ionic– electronic thin films able to generate photocurrents [216, 217]. Apart from Cu2O, there are further types of solid corrosion products that can be formed at the copper–water interface, depending on the composition of the waters. Species more likely to be present in waters are C, CO2, Cl2, Cl , NO2, Fe, Na2SO4, HSO4 , SO42 , S, H2S, HS , H2SO5, S8, and SO2, and these can form copper compounds, such as Cu2S, Cu2O, CuO, CuCl, CuOHCl, Cu(OH)2, CuS, and CuSO5H2O. Thermodynamic data for most of these compounds are given by Bard et al. [218], Garrels and Christ [219], Wagman et al. [220], and others and are being used in the E–pH graphical form [221] to predict the kinetic behavior of the Cu–O–H, Cu–S–O–H, Cu–C–S–O–H, Cu–F–C–S–O–H and other Cu systems that are relevant to aqueous-phase corrosion processes. Distilled water is not very corrosive to copper, but over long periods the water will pick up traces of copper. Copper tanks should be tinned for holding distilled water or for water storage on ships. In soft waters, particularly those containing appreciable amounts of free carbon dioxide, and in carbonated waters in general, the corrosion of copper is much greater. The initial corrosion rate in distilled water may be 0.051–0.16 mm/year (12–7 mdd), but the rate in an aggressive supply water may be as high as 0.26 mm/year (62 mdd). Such waters may be corrosive enough to pick up sufficient copper to form green stains on plumbing fixtures. Hard waters are seldom corrosive to copper because of the protective film of calcium compounds that soon forms. When waters with an appreciable temporary hardness are softened, however, they can become corrosive, especially if heated above 140 C. The calcium compounds necessary to form a protective film are then absent and the sodium bicarbonate formed breaks down with the release of carbon dioxide. Copper is very stable in salt solutions and seawater. When constantly immersed, the corrosion rate of pure copper in

CORROSION IN WATERS

seawater is 0.02–0.07 rnm/year (5–17 mdd) and under varying immersion, as at half-tide, the rate is 0.02–0.1 mm/year (5–25 mdd). In the latter case, arsenical copper is rather more resistant than tough-pitch copper. Under constant immersion the corrosion resistance of copper is two to five times greater than that of mild steel, and under half-tide immersion the advantage of copper is even more notable. Copper is satisfactory in stagnant seawater or where the velocity is < 1 m/s (3 ft/s). The rate of dissolution of copper is, however, large enough to be toxic to marine boring animals, hence its value as a sheathing material for wooden craft and piling. For the same reason, copper compounds are used in marine antifouling paints. Copper is not sufficiently corrosion resistant in rapidly flowing seawater to be suitable for condenser tubes in oceangoing ships or in power stations using tidal water. A number of copper alloys have superior resistance under such conditions (see below). Pitting of copper tubes caused by carbonization of the lubricant used in tube making has been largely eliminated by improving and controlling the tube manufacturing process. Films of manganese oxides, derived by slow deposition from soft moorland waters, can also give rise to localized attack. A beneficial effect, presumably due to the presence of an organic chemical at very low concentration, is observed in the corrosion of copper in certain waters. The organic agent responsible for this action has not been isolated and the way in which it operates is imperfectly understood. E2. Copper Alloys Copper alloys have a long history of successful application in freshwater and saltwater environments. In seawater, in addition to cuprite (Cu2O), Cu2(OH)3Cl forms at the alloy– water interface, independently of alloy composition. Copper alloys (e.g., C44300, C44500, C61300, C68700, C70600, or C71500) are more resistant to corrosion in seawater than pure copper because of the incorporation of relatively corrosionresistant metals such as nickel, or because of the addition of metals such as iron and aluminum that assist in the formation of protective oxide films. Information is summarized below for the copper alloys more commonly used in water. E2.1. Brasses. The common brasses are alloys of copper with 10–50% zinc and often a number of other components including tin, iron, manganese, aluminum, and lead. Zinc dissolves in copper up to 39% to give a single phase alloy, a brass, and with zinc contents in the range of 47–50% the alloy is again single phase, b brass, Between these two, that is with 39–47% zinc, the alloy contains both phases, a þ b brass. In the corrosion of the a brasses the constituents are dissolved according to the copper/zinc ratio, but as the two-phase alloy is approached, appreciably more zinc dissolves. In the duplex

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brasses, dezincification takes place by preferential dissolution of the less noble b phase, and in the pure b alloys only zinc appears in the corrosion products, the copper being redeposited as metal. The stability of brasses in natural waters depends in a complex manner on the dissolved salts, the water hardness, the dissolved gases, and on the formation of protective films. In general, the rates of attack vary from 0.003 to 0.03 mm/year (0.6 to 6 mdd). The corrosion rate is increased by higher concentrations of carbon dioxide and, with the higher zinc brasses, is accompanied by dezincification. Because tin additions are effective in reducing dezincification, naval brass, 63 Cu–36 Zn–1 Sn, is commonly used. For freshwater plumbing, piping of red brass, 85 Cu–15 Zn, is preferred to copper alloys such as Muntz metal, 60 Cu–40 Zn, and leaded yellow brass, 67 Cu–33 Zn, which dezincify in some waters. The corrosion rate of brass by pure water, such as condensate, is very low, 5% tin and with additions of lead and zinc are used. These include the wellknown Admiralty bronze or Admiralty gun-metal, 88 Cu–10 Sn–2 Zn, and naval bronze, 88 Cu–8 Sn–4 Zn. In certain conditions, the corrosion rate of lead–tin bronzes with tin contents > 10% increases to 0.09 mm/year (22 mdd). E2.3. Aluminum Bronze. Aluminum bronzes usually contain not more than 9–10% aluminum and sometimes small additions of manganese and nickel as well. In seawater, they are more corrosion resistant than other copper alloys, with corrosion rate only 10% of that of bronze and 3% of that of brass. The rate of corrosion on an 8.0% aluminum alloy is 0.001–0.003 mm/year (0.3–0.8 mdd) at 30 C (89 F), and 0.01 mm/year (2.3 mdd) at 60 C (143 F). More complex alloys containing up to 11.5% aluminum, 5.5% nickel, 5.0% iron, and 3.5% manganese are casting alloys used in the manufacture of ships’ screws. These alloys are able to withstand the severe demands of service in Arctic waters better than manganese brasses and are much more stable toward erosion and cavitation. Alloys C61300 and C63200 are used in cooling tower hardware in which the makeup water is sewage effluent. Aluminum bronzes resist oxidation and impingement corrosion because of the aluminum in the surface film. E2.4. Silicon Bronzes. Silicon bronzes contain up to 4.5% silicon and sometimes a number of other additions. The 3.0% Si bronzes are stable in natural waters; including seawater, and are very suitable for hot water equipment and for screws used in marine fittings. Corrosion rates for silicon bronzes are similar to those for copper. E2.5. Copper-Nickel Alloys. Alloys based on copper and containing from 5 to 40% nickel have good mechanical propertied at moderately elevated temperatures and excellent resistance to corrosion in many environments and in particular in contact with brackish water and seawater [222–225]. They are more stable than the brasses in flowing water, less susceptible to SCC, and are widely used in shipbuilding for installations involving heat exchange. In this group of alloys,

most attention has been directed to 70 Cu–30 Ni and 90 Cu–10 Ni and the improvement in corrosion resistance under flow conditions obtained by small additions of iron and manganese [226]. In both alloys, the optimum performance, even at flow rates up to 5 m/s (16ft/s), is obtained with an iron concentration of 1% (viz., 69 Cu–30 Ni–1 Fe, and 89 Cu–10 Ni–1 Fe). These alloys are recommended for brackish waters, seawaters, and waters with total dissolved solids >2000 ppm (i.e., 0.2%). They have an advantage in that they do not require cathodic protection as is provided for aluminum brass tubes, 76 Cu–22 Zn–2 Al, by iron anodes bolted to the water boxes or by the boxes themselves. A copper–nickel–iron, 69 Cu–30 Ni–Fe, and also alloys of the Monel type (e.g., 30 Cu–66 Ni with 4% iron and manganese) are used for high-pressure feedwater preheaters. The introduction of copper ions into the boiler is thereby avoided. In recent years, tubes of 69 Cu–30 Ni–1 Fe and 79 Cu–20 Ni–1 Fe alloys have been found to exfoliate—thick layers of corrosion product form and subsequently flake off. Exfoliation is limited to heaters that are operated discontinuously and are used with feedwaters treated with sodium sulfite. Exfoliation may be avoided by using the 89 Cu–10 Ni–1 Fe alloy, which can also be used for equipment evaporating seawater to provide drinking water on board ship. The 90 Cu–10 Ni alloy and the silicon bronzes, with very good mechanical properties and good weldability, are used in the fabrication of storage tanks for hot freshwater. The 69 Cu–30 Ni–1 Fe alloy is used for dealing with condensate and feedwater and has a corrosion rate of < 0.08 mm/year (19 mdd). At 70 C (161 F) under air containing carbon dioxide, the corrosion rate is higher, 0.25–0.47 mm/year (64–106 mdd) at 16-atm pressure. This alloy and the corresponding one containing small amounts of manganese are resistant to cavitation. E3. Microbial Corrosion Copper and copper alloys in general use, with few exceptions, are susceptible to some form of microbial corrosion that usually arises from the activity of a wide range of microorganisms and their metabolic products. In addition, the microorganisms are ubiquitous and are able to colonize surfaces and, by genetic mutation, acquire the ability to adapt to environmental changes. They constitute a dynamic system that is able to change with time. The other important feature associated with colonization of copper surfaces is the subsequent formation of biofilms. Adherence of these is brought about by the release of extracellular polymers. Biofilms up to 100 mm thick are not unusual and in nearly all cases contain trapped bacteria. Such films encourage the growth of these bacteria, resulting in the formation of complex biological systems comprising active bacteria, their metabolites, and the chemical changes generated by the system. Because of this complexity, there are

CORROSION IN WATERS

numerous microbial corrosion processes and mechanisms, although it must be emphasized that these do not involve any new corrosion process. Bacterial colonization under aerated conditions usually results in the formation of differential aeration and concentration cells due to the uptake of oxygen by the microbial colony. The oxygen concentration under these colonies becomes depleted and localized pitting corrosion can take place. In some cases, particularly when the iron-oxidizing bacteria are involved, corrosion tubercles develop. Such organisms also facilitate the accumulation of chlorides, resulting deposition of ferric and manganese chlorides. Tuberculation, like slime films, provides a suitable habitat for the sulfate-reducing bacteria (SRB), which can stimulate corrosion by a number of mechanisms, but, under these conditions, it is likely by the releasing of sulfide ions. Because copper and copper alloys rely on the presence of a stable oxide film for their corrosion resistance, they are susceptible to corrosion should the film be damaged or the oxygen shielded from the metal by the biofilm. Under certain circumstances, corrosion can be stimulated by the chemistry within the biofilm and the extracellular polymers that are generated. Chelation of copper ions contributes to the formation of galvanic cells. The mechanism of corrosion may be controlled by the potential difference created between cuprous and cupric ions in the biofilm. In addition to anodic metabolites, some bacteria are able to generate other substances, such as carbon dioxide and ammonia, the latter being a cause of SCC of copper and its alloys. All these environments are of concern, but none perhaps is as important as those in which SRB are active. These prolific organisms cause considerable corrosion damage and pollution of the environment. The mechanism of corrosion involving these bacteria is complex and not completely understood. Nevertheless, it is important to emphasize that at present five hypotheses exist to explain the involvement of SRB in the corrosion process. These include (1) the depolarization of the cathode and utilization of molecular hydrogen, (2) corrosion by sulfide ions, (3) galvanic corrosion due to the formation of iron sulfide films, (4) corrosion due to the formation of elemental sulfur, and (5) the production of a corrosive volatile phosphorus compound. Irrespective of these mechanisms, most copper alloys are susceptible to SRB activity. The rate of corrosion that can arise from SRB activity can be very high particularly when it is controlled by the action of sulfur and its compounds. Digenite (Cu5S9), chalcocite (Cu2S), and covellite (CuS) have been produced by SRB on copper surfaces [227, 228]. Macdonald et al. [228] reported the formation of djurleite (Cu1.96S) from SRB activity on copper alloys. Chalcocite is the most characteristic corrosion product in SRB-induced

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corrosion of copper. Baas-Becking and Moore [227] reported that chalcocite cannot be formed abiotically at room temperature and that microbiological formation of chalcocite is a product of digenite, the first species formed during SRBinduced corrosion of copper alloys. Djurleite is important in SRB-induced corrosion of copper alloys because it forms a protective sulfide film [229] and is difficult to synthesize abiotically at room temperature [230]. McNeil and Little [231] analyzed sulfide mineral deposits on copper alloys colonized by SRB in order to identify specific mineralogies that could be used to fingerprint SRB activity. The copper sulfide mineral found in all combinations of copper-containing substrata and cultures was chalcocite. The authors concluded that the presence of chalcocite was an indicator of SRB-induced corrosion of copper. The compound was not observed in sterile controls and its presence in near-surface environments could not be explained thermodynamically. A range of well-defined types of pitting corrosion of copper were reviewed by Mattsson [40] but none of these was considered to be associated with microbial activity. Since then, two new forms of pitting have been reported that have microbial origins. One of these, termed “pepper-pot pitting” occurred in large institutional buildings in southwestern Scotland, UK [232], whereas the other, frequently termed type l12 pitting, has been observed in tubes from Saudi Arabia, Germany, and England [233]. Examination of affected copper tubes from buildings in Glasgow showed that there was some superficial corrosion due to deposit attack (differential aeration), but that perforation was at sites containing pepper-pot pits. In cross section the pits had a conical gray cap, often hollow, of corrosion products consisting mainly of copper sulfate and cupric oxide. Removing this cap revealed the pepper-pot cluster of pits. Pit morphology (i.e., number, size, and shape of pits) varied slightly within a group. In some, but not all, cases, sulfides could be positively detected. The pits were located in a membrane of cuprous oxide or copper and beneath that were crystals of cuprous oxide extending from the perforations outward in a hemispherical fashion. Further examination of these tube bores has shown that a biofilm can always be lifted from the surface and stained using the periodic acidSchiff stain [234], showing the presence of polysaccharides (by-products of bacteria), which form in copious quantities when biofilms develop. This type of pitting resembles the classical type 3 of pitting [40]. The other new type of pitting occurs in cold, warm, and hot water systems and exhibits features of both type 1 and type 2 pitting (Fig. 56.6) [235]. It resembles type 1 pitting in that the pits are hemispherical and contain soft crystalline cuprous oxide with varying amounts of cuprous chloride under a cuprous oxide membrane. It resembles type 2 pitting in that the oxide on the surface between the

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FIGURE 56.6. Unusual type of pitting corrosion. (Reproduced with permission from Wiley-VCH [235].)

pits is largely cupric. The mounds above the pits are principally basic copper sulfate, often with a deposit of powdery cupric oxide around the periphery and on parts of the deposit itself. An interdisciplinary approach for studying type l12 pitting of copper in potable water systems is reported by Chamberlain et al. [236]. Their electrochemical results are explained on the basis of diffusion processes and multiple Donnan equilibria under the assumption of ion selectivities of mainly anionic exopolymers [237]. Optimizing the alkalinity of potable water and using different tempers and surface conditions are the most promising and economically feasible methods to control type l12 pitting. Bicarbonate dosing, already used to control several corrosion problems, is beneficial only within a certain optimum concentration range of bicarbonate, as indicated by practical experience and results from limited laboratory experiments. It is also well known (DIN 50930, part 5) [238] that in cold water pitting corrosion damage is enhanced in hard water with high alkalinity, and also at elevated temperature [>40 C (> 107 F)] in soft water with low alkalinity. Any measure to improve the performance of different tempers and surface conditions of copper in contact with potable water must not have an adverse effect on tubes already installed and covered by protective layers [238]. In addition to those cases described above, there have been other reports of pitting corrosion in copper and copper alloys exposed to potable water and other industrial water systems; references [35, 107, 239–241] give further information. E4. Biofouling Slime algae, sea mosses, sea anemones, barnacles, oysters, and mussels attach themselves to marine structures, such as pilings and offshore platforms, boat hulls and even the insides of pipes and condensers. This phenomenon, called marine biofouling, is common in open waters an estuaries. Either mechanical removal or prevention (e.g., by renewable antifouling coatings) is often required.

Marine organisms adhere to some metals and alloys more readily than they do to others. Copper-based alloys have very good resistance to biofouling, and this property can be used to good advantage. Copper–nickel is used to minimize biofouling on intake screens, seawater pipework, water boxes, mesh cages in fish farming, marine craft, and offshore structures. A prime example of this was demonstrated in 1987 when two early solid copper–nickel hulled vessels, the Asperida II and Copper Mariner, were inspected after 21 and 16 years, respectively, of service. Neither vessel had required hull cleaning or suffered significant hull corrosion during that time. Alternative antifouling methods for hulls would have required recoating after 30 months for copper-based paints or 5 years for organotin copolymer coatings. In the short term, coatings are the less expensive alternative, but based on life-cycle costing, the copper–nickel hull is more cost effective. Cost studies to date have estimated the payback time to be 4–7 years for commercial vessels. With the current environmental concerns and restricted use of tributyltin copolymer paints, the use of copper–nickel to control corrosion and biofouling of ship hulls should be considered. It is also timely to review products incorporating copper–nickel that can be used for protecting marine structures from biofouling. The 90 Cu–10 Ni alloys were developed from the very successful 70 Cu–30 Ni alloys as a result of conservation of nickel as a strategic metal during World War II [242]. About 1.5% iron increased resistance to erosion–corrosion, but additions >3.5% increase susceptibility to deposit attack and biofouling [243]. Copper–nickel alloys have traditionally been used for pipe work carrying seawater as they exhibit both corrosion and biofouling resistance, but more recently the 90 Cu–10 Ni alloys have been used in marine construction, such as intake screens for seawater-cooled power plants, fish-rearing cages, sheathing for oil and gas platforms, and boats [244]. A number of small boats have been constructed from copper– nickel alloys since the 1940s and have performed well, although at rather low speeds. Trials with a ship travelling

CORROSION IN WATERS

at higher velocities are encouraging, especially as the hull material is self-smoothing [244]. Average corrosion rates reported are about 0.02 mm/year. (0.8 mpy, 5 mdd). Although 90 Cu–10 Ni alloys have very good corrosion resistance, at least at flow velocities below 4 m/s, they do corrode slightly. This slow release of copper is responsible for antifouling properties, and if it is halted by a cathodic protection system or galvanic coupling to a less noble metal, then biofouling will occur. Antifouling properties of 90 Cu–10 Ni depend on the environmental conditions at the exposure site. Deployment in open water with moderate to considerable wave action is likely to result in an essentially clean surface, but exposure in low-water-velocity harbors and basins may result in the intermittent development of thick, complex fouling communities. Although macrofouling is often absent, microfouling usually occurs and is not restricted to prokaryotes but can include unicellular algae, protozoa, and other eukaryotes. The occurrence of microbial communities on copper alloys has been discussed in detail [222, 242, 245–249]. Efird [250] recorded that “heavy sliming” was present after 12 months exposure in the sea for a range of copper alloys. The development of microbially dominated biofilms on copper–nickel alloys has not really been appreciated by many engineers and corrosion scientists, largely because thin biological films are inconsequential when calculating, for instance, the increased drag and wave loading due to macrofouling or perhaps the increase in weight due to shellbearing animals. Moreover, it is generally believed that copper ions are lethal to the majority of organisms. Early investigations described microbially influenced corrosion (MIC) of copper alloys [251, 252]. Bacteria isolated on 70 Cu–30 Ni showed considerable copper tolerance, up to 2000 ppm copper [252]. These tests were carried out in suspension culture, and it is now well established that microorganisms residing within a mucilaginous biofilm may show much higher levels of tolerance to toxic compounds, antibiotics and the like. This research showed that MIC occurs on copper–nickel condenser tubes. Mechanisms proposed for MIC of copper–nickel included formation of corrosive metabolic products (e.g., ammonia, organic and inorganic acids, and hydrogen sulfide) that cause cathodic depolarization or catalysis of other corrosion reactions, and anaerobic sulfide production [252]. The effects of sulfide to cause corrosion can be seen long after the source of sulfide has disappeared. Doping a passivating cuprous oxide corrosion product layer by sulfide may lead to very high rates of corrosion once the material is exposed to aerated conditions [253, 254]. E5. Condensers and Heat Exchangers Apart from the wide range of waters that must be used for cooling and the consequent variability in corrosion

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characteristics, some of the most corrosive conditions are encountered in power stations and ships that draw their cooling water from estuaries and harbors. These waters are polluted to varying extents depending on location and the time of year. Thus water drawn from an area near the outfall of a sewerage works will be heavily polluted and will cause localized corrosion of condenser tubes. Corrosion initiated by polluted water, which is sometimes intergranular, may often continue even after changing to a clean water. The concentration of pollutants (e.g., hydrogen sulfide, which arises from the decomposition of organic matter), is likely to be greatest in late summer. Oxygen is consumed by organic matter, and SRB become active. Hydrogen sulfide is probably the most powerful pollutant of waters and can cause rapid perforation of materials (e.g., brass). Another important pollutant, cystine, is produced by the breakdown of organic material, such as seaweed, and is present in both offshore and in-shore waters, the concentration being particularly large in harbors and estuaries. Cystine is a very efficient cathodic depolarizer and can produce intense attack even in the absence of oxygen. If cystine is present in sufficient concentration, however, a protective film forms on copper alloys consisting of a copper–cystine complex. This film may break down or blister if it becomes too thick (e.g. under impingement) and results in the highly corrosive situation of a small anode and large cathode. In the construction of condensers and heat exchangers, the metals and alloys used are copper, brasses, bronzes, and cupronickel alloys [255]. Among these alloys C44300, C44400, C44500, C68700, C613O0, C63200, C63000, C95400, C95500, C95800, C70600, C71500, and C12000 provide satisfactory and economical performance for condenser and heat exchanger tubing. The corrosion resistance of these latter materials depends both on the inherent nobility conferred on them by copper and nickel and protection by the layer or film of corrosion product. The properties of this corrosion product, its continuity, adhesion to the metal surface and its ability to form, to be maintained, and to survive the action of erosive forces of highvelocity seawater, determine the alloy chosen for a given use. The effects of velocity are complicated by the presence in water of entrained air bubbles and abrasive materials such as sand, which can cause erosion and impingement attack leading to film breakdown and accelerated corrosion. The incidence of impingement attack on copper–alloy tubes is reduced when entrained air is removed from the cooling water. In other cases, dissolved and entrained air in the water may help to form and maintain protective films on the more resistant materials. Iron corrosion products also help certain alloys to form protective films. In one case, introduction of ferrous sulfate into the cooling water at a power station reduced the attack on aluminum brass [256].

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Pure deoxidized copper is limited to applications where purity of metal is essential. It has good resistance to practically all types of freshwater, but it is not serviceable in salt water and waters polluted with sulfur compounds. [Above 200 C (395 F) there is a significant loss of strength.] Impingement attack occurs when the velocity of the water is high and no protective scale forms (e.g., in soft water high in free carbon dioxide). Where this type of attack is expected, reinforcement or protection are necessary or a different material should be selected. Arsenical copper containing 0.15–0.5% arsenic has a higher strength and fatigue limit but retains ductility. Its main use is in condensers and heat exchangers for use in freshwater. It is not recommended for waters containing hydrogen sulfide and other sulfur compounds, acid mine water, and salt or brackish waters. Copper is relatively sensitive to metal ion concentration cells, and so the velocity of liquid across the surface is an important factor in controlling the corrosion rate. The resistance of copper to corrosion by rapidly flowing seawater is unsatisfactory, and it is not suitable for condenser tubes in ocean-going ships or in power stations using tidal waters. The use of Muntz metal, 60 Cu–40 Zn, is mainly confined to steam condensers operating at low temperatures and using freshwaters from rivers, lakes, and wells. However, it has good resistance to hydrogen sulfide and other sulfur compounds and, with additions of lead, is commonly used in heat exchangers cooled with seawater. Red brass is also serviceable in heat exchangers using freshwaters and is not susceptible to either SCC or dezincification, but is likely to be severely corroded in the presence of sulfur compounds. The addition to brass of small amounts of alloying elements is beneficial. Naval brass, 60 Cu–39 Zn–1 Sn, and Admiralty brass, 70 Cu–29 Zn–1 Sn, are used to resist seawater in exchangers if the flow rate is 125 ppm calcium carbonate) form adequate protective films on lead and therefore attack is negligible. The presence of salts such as silicates increase the hardness and film thickness. Nitrates interfere with protective film formation, causing increased corrosion. In soft aerated waters, corrosion depends on both hardness and oxygen content. With hardness levels of 0.015 ppm of lead is generally considered contaminated and unsafe for consumption. Such contamination can be caused even in situations where from a service point of view the corrosion rate of lead is negligible, and can include the use of lead solders in copper pipe. Corrosion can be accelerated if carbonic acid is present in water, converting calcium carbonate deposits into soluble calcium bicarbonate. Also the presence of organic acids whose lead salts are soluble can promote corrosion. Filmforming lime or sodium silicate can be added to water to lower corrosion rate.

TABLE 57.2.

B4. Demineralized Water The corrosion rates of chemical lead, Pb–6% Sb and Pb–2% Sn in demineralized water are shown in Table 57.3. These rates were measured after 21 days of immersion. Corrosion rates are observed to be very low; however, the chemical lead did show some waterline attack.

Corrosion of Chemical Lead in Industrial and Domestic Watersa Temperature F

C

Aeration

Agitation

mddb

mpyc

70–100 68 67 72 60–85

21–38 20 19 22 16–29

None Yes Yes Yes Complete

Slow Slow Slow Slow None

6.75 2.08 2.20 1.98 41.7

0.85 0.26 0.28 0.25 5.3

0.15 m/s (0.5 ft/s)

2.95

0.38

2.9

0.37

Kind of Water Condensed steam, traces of acid Mine water, pH 8.3,110 ppm hardness Mine water, 160 ppm hardness Mine water, 110 ppm hardness Cooling tower, oxygenated Lake Erie water Los Angeles aqueduct water, treated by chlorination and copper sulfate Spray cooling water, chromate treated a

See [1]. In milligrams/square decimeter/day (mg/dm2/day). c In mils/year. b

Corrosion Rate

Ambient 60

16

Yes

790

LEAD AND LEAD ALLOYS

may be pollution, rate of flow, wave action, sand or silt content, temperature, and marine growth. Insufficient data have been produced to reveal the effects of these factors. A summary of the effects of alloying of lead on corrosion rate is shown in Table 57.4.

C. CORROSION IN CHEMICAL MEDIA C1. Effect of pH on Corrosion As illustrated in Figure 57.2, the corrosion of lead is accelerated by both acids and alkalis because of its amphoteric nature. Significant corrosion can occur < pH 5 or > pH 10. C2. Inorganic Acids The acids to which lead offered good resistance were listed earlier. Lead of 3 mm (18 in) of thickness or greater usually gives years of service in equipment handling these chemicals. High acid velocities may increase the rate of corrosion. Figure 57.3 shows the effect of velocity of sulfuric acid across a lead base at 25 C. Lead satisfactorily resists all but the most dilute strengths of sulfuric acid. It performs well with concentrations up to 95% at ambient temperatures, up to 85% at 220 C, and up to 93% at 150 C. Below a concentration of 5%, corrosion rates increase but are still relatively low. At lower concentrations, Pb–Sb alloys are recommended. Similar behavior patterns occur with chromic, sulfurous, and phosphoric acids. The behavior of lead in sulfuric acid is shown in Figure 57.4 and its behavior with air-free hydrofluoric acid in Figure 57.5. Tabular data for the behavior of lead with phosphoric, hydrochloric, and a hydrochloric–ferric–chloride mixture is tabulated in Tables 57.5–57.7. Most concentrations of nitric, acetic, and formic acids corrode lead at a rate high enough to prevent its use; however, nitric acids of 52–70% do not

FIGURE 57.1. Effect of oxygen on corrosion of lead submerged in distilled water at 25 C. [Reprinted with permission from H. H. Uhlig (Ed.), Corrosion Handbook, Wiley, New York, 1948.]

TABLE 57.3. Corrosion of Lead in Demineralized Watera Lead

mpy

mdd

Chemical lead (ASTM B29) 6% Antimonial lead 2% Tin–lead

2.3 0.2 0.6

18 1.6 4.8

a

See [1].

B5. Seawater The corrosion of lead in seawater is relatively low and may be slowed by encrustations of lead salts. Seawater normally has a salt content of 3–4%. Other factors influencing corrosion TABLE 57.4. Corrosion of Lead and Lead Alloys in Seawatera Metal

Days

Depth m

Chemical lead (ASTM B29)

181 197 123 181 197 123 181 197 123 181 197 123

1.5 713 1719 1.5 713 1719 1.5 713 1719 1.5 713 1719

Pb–0.04% Te

Pb–6% Sb

Pb–33% Sn

a

See [2].

Thickness loss mpy

Weight loss mdd

Pit Depth

1.2 0.3 0.8 1.0 0.3 1.1 1.2 0.3 0.8 3.7 0.5 0.5

9.4 2.4 6.3 7.9 2.4 8.7 9.4 2.4 6.3 29.1 3.9 3.9

0 0 0 0 0 0 0 0 0 0 0 0

Remarks Uniform attack Uniform attack Uniform attack Uniform attack Uniform attack Uniform attack Uniform attack Uniform attack Uniform attack Uniform attack Uniform attack Uniform attack

CORROSION IN CHEMICAL MEDIA

791

FIGURE 57.3. Effect of velocity across metal face on the corrosion of lead in 20% sulfuric acid at 25 C. (Adapted from [1], used with permission.)

FIGURE 57.2. Corrosion of lead in contact with various acid and alkali solutions. (Reproduced with permission from [2].)

produce significant attack, while dilute concentrations give rapid attack. This pattern is also true for hydrofluoric, acetic, and sodium sulfate acids. Adding sulfuric acid to acids corrosive to lead will often lower the corrosion rate. For example, even though nitric acid of 1275 mm/year (50 mils/year)

Negligible corrosion, lead recommended for use. Practically resistant, lead recommended for use. Lead may be used for this effect on service life can be tolerated. Corrosion rate too high to merit any consideration of lead.

C4. Alkalis The action of strong alkalis on lead is not as rapid as by acids like hydrochloric or acetic but is greater than the action of alkalis on iron. For certain purposes, however, corrosion of lead in contact with NaOH or KOH tip to 30% concentration at 25 C and up to 10% concentration at higher temperatures (90 C) is tolerable. For example, lead has proven useful in the refining of petroleum oil, where a sulfuric acid treatment is followed by an alkaline solution treatment in the same lead-lined tank. Although Ca(OH)2 (lime) solutions are saturated at 0.1% Ca(OH)2 at 25 C, they have been found to corrode lead severely when trickling over the surface. This can occur with fresh concrete or with long time seepage of water through cured concrete. In such cases, dissolved oxygen appears to be necessary for corrosion.

C6. Salts Lead is generally resistant to corresponding salts of the acids to which it is resistant Nitrates and alkaline salts tend to be actively corrosive. C7. Gases Lead resists the actions of chlorine, wet or dry, up to 100 C, but that of bromine only when dry and at lower temperatures. Sulfur dioxide and trioxide, wet or dry, are frequently handled in lead pipes. Hydrogen sulfide, with or without moisture, can be handled with lead providing erosion at high velocities is avoided. Hydrogen fluoride is actively corrosive, therefore, lead is not recommended as a container for this gas.

794


TABLE 57.12. Corrosion Rate of Lead in Chemical Environmentsa Chemical Abietic acid Acetaldehyde Acetaldehyde Acetanilide Acetic acid Acetic acid Acetic anhydride Acetoacetic acid Acetone Acetone cyanohydrin Acetophenetidine Acetophenone Acetotoluidide Acetyl acetone Acetyl chloride Acetyl thiophene Acetylene, dry Acetylene tetrachloride Acridine Acrolein Acrylonitrile Adipic acid Alcohol, ethyl Alcohol, methyl Alkanesulfonic acid Alkyl aryl sulfonates Alkyl naphthalene sulfonic acid Allyl alcohol Allyl chloride Allyl sulfide Aluminum acetate Aluminum chlorate Aluminum chloride Aluminum ethylate Aluminum fluoride Aluminum fluorosulfate Aluminum fluosilicate Aluminum formate Aluminum formate Aluminum hydroxide Aluminum nitrate Aluminum potassium sulfate Aluminum potassium sulfate Aluminum sodium sulfate Aluminum sulfate Aminoazobenzene Aminobenzene sulfonic acid Aminobenzoic acid Ammophenol Aminosalicylic acid Ammonia Ammonium acetate Ammonium azide Ammonium bicarbonate

Temp. ( C)

Concentration (%)b

Corrosion Class

24 24 24–100 24 24 24 24 24 24–100 24–100 24 24–100 24 24–100 24 24–100 24 21 24–106 24–106 24–100 24–100 24–100 24–100 24 24–100 93 24 24 24 24–100 24–100 24 24–100 24–100 24 24 24 100 24–100 24 24–100 24–100 24–100 24–100 24 24–100 24–93 24–100 100–149 24–100 25 24 24–100

— — — —

D A B A B C A B A B B B B B A B A B B B A A A A D B C B C D A B B B B A B B D B B A B B A C B B B C B B B B

Glacial — — — 10–90 — — — — — — — — Liquid 10 10 — — 10–100 10–100 — — — — — — 10–20 — 0–10 — 10–20 15 — — — 10 10 10–20 20–100 10 — — — — — — 10–30 3.85 — 10


TABLE 57.12.

795

(Continued )

Chemical Ammonium bifluoride Ammonium bisulfate Ammonium carbamate Ammonium carbonate Ammonium chloride Ammonium citrate Ammonium diphosphate Ammonium fluoride Ammonium fluosilicate Ammonium formate Ammonium hydroxide Ammonium hydroxylamine Ammonium metaphosphate Ammonium nitrate Ammonium oxalate Ammonium persulfate Ammonium phosphate Ammonium picrate Ammonium polysulfide Ammonium sulfamate Ammonium sulfate Ammonium sulfide Ammonium sulfite Ammonium thiocyanate þ NH4OH Ammonium tungstate Amyl acetate Amyl chloride Amyl laurate Amyl phenol Amyl propionate Aniline Aniline hydrochloride Aniline sulfate Aniline sulfite Anthracene Anthraquinone Anthraquinone sulfonic acid Antimony chloride Antimony pentachloride Antimony sulfate Antimony trifluoride Arabic acid Arachidic acid Arsenic acid Arsenic trichloride Arsenic trioxide Ascorbic acid Azobenzene Barium carbonate Barium chlorate Barium chloride Barium cyanide Barium hydroxide Barium nitrate

Temp. ( C) 24 24–100 24–149 24–100 24 100 24–100 24 24–52 100 27 20–100 24 24–100 24 24–100 66 24–100 24–100 24–100 24 24–100 24–100 24 24 24 24 24–100 200 24–100 20 24 24–100 24–100 24–100 24–100 24–100 24 24–100 100 24–100 24–100 24 24 24–100 24–100 24 24–100 24 24–100 24–100 24 24 24–100

Concentration (%)b 10 — — 10 0–10 — 10 0–20 20 10 3.5–40 34 10 10–30 10–30 10–30 — 10 10 10 — 10 10–40 — 10 80–100 — — — — — 10 — — — — 10–30 — 90–100 — 50–70 — — 10 — — — — — 20 10 10–70 10 10–30

Corrosion Class B A A B B D B B B C A B B D D B A B B B B C B A D B D B D B A D B B B B B C B C A B B B B B D B D B B D D B (continued )

796


TABLE 57.12. (Continued ) Chemical Barium peroxide Barium polysulfide Barium sulfate Barium sulfide Benzaldehyde Benzaldehyde sulfonic acid Benzamide Benzanthrone Benzene Benzene hexachloride Benzene sulfonic acid Benzene sulfonic acid Benzidine Benzidine disulfonic acid 2.2 Benzidine 3 sulfonic acid Benzilic acid Benzobenzoic acid Benzocathecol Benzoic acid Benzol Benzonitrile Benzophenone Benzotrichloride Benzotrifluoride Benzoyl chloride Benzoyl peroxide Benzyl acetate Benzyl alcohol Benzylbutyl phithalate Benzyl cellulose Benzyl chloride Benzyl ethyl aniline Benzylphenol Benzylphenol salicylate Benzylsulfonilic acid Beryllium chloride Beryllium fluoride Beryllium sulfate Boric acid Bornyl acetate Bornyl chloride Bornyl formate Boron trichloride Boron trifluoride Bromic acid Bromine Bromobenzene Bromoform Butane Butanediols Butyl acetate Butyl benzoate Butyl butyrate Butyl glycolate

Temp. ( C) 24 100 24–100 24 24 24–100 24–100 24–100 24 24 24 100 100 24–100 24–100 24–100 24–100 24–100 24 24 24–100 24–100 24–100 24–100 100 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24–100 100 24–100 24–100 24–149 24–100 24–100 24–100 24–100 24–204 24–100 24 24–100 24–100 24 24 24 24–100 24–100 24–100

Concentration (%)b 10 — — 10 10–100 — — — — — 10–100 — — — — 10–100 — — — 100 — — — — — — — — — — — — — — — — — 10–50 10–100 — — — — — 10–30 — — — — — — — — —

Corrosion Class D D B B D B B B B B B D B B B B B B D A A A B B C B B B B B B B B B B D B B B B B B B A B B B B A B B B B B


TABLE 57.12.

797

(Continued )

Chemical Butyl mercaptan Butyl oxalate Butyl phenols Butyl phthalates Butyl stearate Butyl urethane Butyric acid Butyric aldehydes Butyrolactone Cadmium cyanide Cadmium sulfate Calcium acetate Calcium acid phosphate Calcium benzoate Calcium bicarbonate Calcium bisulfite Calcium bromide Calcium carbonate Calcium chlorate Calcium chloride Calcium chromate Calcium dihydrogen sulfite þ SO2 Calcium disulfide Calcium fluoride Calcium gluconate Calcium hydroxide Calcium lactate Calcium nitrate Calcium oxalate Calcium phosphate Calcium pyridine sulfonate þ H2SO4 Calcium stearate Calcium sulfaminate Calcium sulfate Calcium sulfide Calcium sulfite Camphene Camphor Camphor sulfonic acid Capric acid Caprolactone Capronaldehyde Capronaldehyde Carbazole Carbitol Carbon disulfide Carbon fluorides Carbon tetrabromide Carbon tetrachloride (dry) Carbonic acid Carnallite Carotene Cellosolves Cellulose acetate

Temp. ( C)

Concentration (%)b

24 24 24 24–100 24–100 24–100 24 24–100 24–100 24 24–100 24 24 24–100 24 24 24–100 24 24 24 24–100 24 24 24–100 24–100 24 100 24 24–100 100 24 24–100 24–100 24–100 100 24–100 24–100 100 24 24–100 24–100 24 52–100 25–100 25–100 25–100 25–100 100

— — — — — —

b

24 24–100 24–100 24–100 24

10–100 — — — 10–30 20 10–30 — — — 30 20 10–30 20 10 5 — — — 10 10 10 10 10 20 10 — 10 — 10 — — 20–100 — — — — — — — — — 100 — — — — —

Corrosion Class C B C B B B D B B D A B B B C B B D B A B A B B B D B D B B A B A B C B B A C B B A B B B A B C A D A A A A (continued )

798


TABLE 57.12. (Continued ) Chemical Cellulose acetobutyrate Cellulose nitrate Cellulose tripropionate Cerium fluoride Cerium sulfate Cesium chloride Cesium hydroxide Cetyl alcohol Cetyl alcohol Chloroacetic acid Chloral Chloramine Chloranil Chloranthraquinone Chlordane Chlorethane sulfonic acid Chloric acid Chlorine Chlorine dioxide Chloroacetaldehyde Chloroacetone Chloroacetyl chloride Chloro–alkyl ethers Chloroaminobenzoic acid Chloroaniline Chlorobenzene þ SO2 Chlorobenzotrifluoride Chlorobenzoyl chloride Chlorobromomethane Chlorobromopropane Chlorobutane Chloroethylbenzene Chloroform Chlorohydrin Chloromethonic ester Chloronaphthalene Chloronitrobenzene Chlorophenohydroxy acetic acid Chlorophenol Chloroquinine Chlorosilanes Chlorosulfonic acid Chlorosulfomc acid þ 50% SO3 Chlorotoluene Chlorotoluene sulfonic acid Chlorotoluidine Chlorotrifluoro ethylene Chloroxylenols Chloroxylols Cholesterol Chromic acid Chromic chloride Chromic fluoride Chromic hydroxide

Temp. ( C)

Concentration (%)b

Corrosion Class

24–100 24–100 24–100 24–100 100 24–100 24 24 100 24 24–100 24 24–100 24–100 24–100 100 24 37 6 24 24–100 24 24–100 24–100 24–100 18 24–100 24–100 24 24–100 24 24–100 24–BPc 24–100 24–100 24–100 24–100 24–100 24 24 24–100 24 14 24–100 24 24–100 24–100 24 24–100 24–100 24 24–100 24–100 24–100

— — — — — —

B B B B C B D B C B B B B B B C D B B B B B B B B A B B B B B B B B B B B B C C B C C B C B B C B B B B B B

10 — — — — 10–30 — — — — 10 — Gas — — — — — — — — — — — — — — — — — — — — — — — 40 — — — — — — — — — — —


TABLE 57.12.

799

(Continued )

Chemical Chromic phosphate Chromic sulfate Chromium potassium sulfate Chromium sulfate (basic) Chromyl chlorides Critic acid Citric acid Cobalt sulfate Copper chloride Copper sulfate m-cresol þ 10% water m-cresol þ 10% water o-cresol þ 10% water o-cresol þ 10% water Cresote Cresylic acid Cresylic acid Crotonaldehyde Crotonic acid Cumaldehyde Cumene Cumene hydroperoxide Cyanamide Cyanoacetic acid Cyanogen gas Cyclohexane Cyclohexanol Cyclohexanol esters Cyclohexanone Cyclohexene Cyclohexylamine Cyclopentane DDT Dialkyl sulfates Dibenzyl Dibutyl phthalate Dibutyl thioglycolate Dibutyl thiourea Dichlorobenzene Dichlorodifluoro-methane–(Freon 12) Dichlorodiphenyldichloroethane (DDD) Dichloroethylene Diethanolamine Diethyl ether Diethylamine Diethylaniline Dielhylene glycol Difluoroethane Diglycolic acid Dihydroxydiphenylsulfone Diisobutyl Dimethyl ether Dioxane Diphenyl

Temp. ( C) 24–100 24–100 24–100 24–100 24–100 24–80 24 24 24 24–100 25 BPc 25 BPc 24 24 24 24–100 24 24–100 24–100 24–100 24–100 24 24 24 24 24–100 24 24–100 24 24–100 24 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24 24 24 24–100 24–52 24–100 24 24–100 24–100 24–100 24–100 24–100

Concentration (%)b — 10 10 20–50 — 10–30 50–100 10–30 10–40 10–70 Liquid Vapor Liquid Vapor 90 90 100 — — — — — — — — — — — — — — — — — — — — — 10–100 90 — — — — — — — — — — — — — —

Corrosion Class B B B B B B D B D B B D B D D D B B D B B D B D D B B B B B D B B B B B B B B A B A B B D B B B D B B B B B (continued )

800


TABLE 57.12. (Continued ) Chemical Diphenyl chloride Diphenylamine Diphenylene oxide Diphenylpropane Epichlorohydrin Ethane Ether Ethyl acetate Ethyl benzene Ethyl butyrate Ethyl cellulose Ethyl chloride Ethyl ether Ethyl formate Ethyl lactate Ethyl mercaptan Ethyl stearate Ethyl sulfonic acid Ethyl sulfonic acid Ethylene Ethylene bromide Ethylene chlorohydrin Ethylene chlorohydrin Ethylene cyanohydrin Ethylene cyanohydrin Ethylene dibromide Ethylene dichloride Ethylene glycol Ethylene oxide 2-Ethylhexoic acid Ferric ammonium sulfate Ferric chloride Ferric ferrocyanide Ferric sulfate Ferrous ammonium sulfate Ferrous chloride Ferrous sulfate Fluoboric acid Fluocarboxylic acid Fluorine Fluosilicic acid Formaldehyde Formamide Formic acid Furfural Gluconic acid Glutamic acid Glycerol Glycerophosphoric acid Glycol monoether Glycolic acid Glycolic acid Heptachlorobutene Heptane

Temp. ( C)

Concentration (%)b

Corrosion Class

24–100 24–100 24–100 24–100 24 24–100 24 24–80 24–100 24–100 24–100 24–100 24–100 100 24–100 100 24–100 24 100 24–100 100 24 52 24 52–100 24 25–100 –7 24 71 24–100 24 66–93 25–80 24 24 24–100 24 24 24–100 45 24–52 24–100 24–100 24–100 24 24 24 24 24–100 24 100 24 24–100

— — — — — — — — — — — — — — — — — — — — —

B A B B A A B B B B B B B C B D B B C A B A B A B D B B B C A D A A B C B C D A D B B D B B D B B B B D B A

90 100 — 100 90 — 50 — 96 10–20 20–30 — 10–20 10 10–30 10 30 — — 10 20–100 — 10–100 — 10–100 — — — — 10–100 10 — —


TABLE 57.12.

801

(Continued )

Chemical Hexachlorobutadiene Hexachlorobutene Hexachloroethane Hexamethylene tetramine Hydrazine Hydriodic acid Hydrobromic acid Hydrochloric acid Hydrofluoric acid Hydrogen bromide (Anh HBr) Hydrogen chloride (Anh HCl) Hydrogen peroxide Hydrogen sulfide Hydroquinine Hydroxyacetic acid Hypochlorous acid Iodine Iodoform Isobutyl chloride Isobutyl phosphate Isopropanol Lactic acid Lead acetate Lead arsenate Lead azide Lead chloride Lead chromate Lead dioxide Lead nitrate Lead oxide Lead peroxide Lead sulfate Lithium chloride Lithium hydroxide Lithium hypochlorite Lithopone Magnesium carbonate Magnesium chloride Magnesium chloride Magnesium hydroxide Magnesium sulfate Maleic anhydride Malic acid Mercuric chloride Mercuric sulfate Mercurous nitrate Mercury Methanol Methyl ethyl ketone Methyl isobutyl ketone Methylene chloride Monochloroacetic acid Monochlorobenzene Monoethanolamine

Temp. ( C)

Concentration (%)b

24–100 24 24–100 24–100 24–100 24–100 24 24 24 100 24 24 24 24–100 24 24 24 24–100 24 24 24 24 24 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24 24–80 24 24 24 24 24 24–100 27 100 24 24–100 24 24 30 24–100 24–100 24–100 24 24 171

— — — 10–40 20–100 10–50 10–70 0–10 2–10 — 100 10–30 90–100 10 — — — 10 — — — 10–100 10–30 — — — — — — — — — 10 — 10 — 10 0–10 10–100 10–30 10–60 10 — 10 10 — 100 — 10–100 10–100 — 20–100 90 —

Corrosion Class A B B B D D D C B D A D B B A D D B B B A D D B B B B B B B B B B D A A D C D D B C B C B D D B B B B D D C (continued )

802


TABLE 57.12. (Continued ) Chemical Naphthalene Naphthalene sulfonic acid þ H2SO4 Nickel ammonium sulfate Nickel nitrate Nickel sulfate Nitric acid Nitrobenzene Nitrocellulose Nitrochlorobenzene Nitroglycerine Nitrophenol Nitrosyl chloride Nitrosylsulfuric acid Nitrotoluene Nitrous acid Oleic acid Oxalic acid Oxalic acid þ 1.5–3% H2SO4 Pentachlorethane Perchloroethylene Persulfuric acid Phenol Phenolsulfonic acid Phenyl isocyanate Phosgene Phosphoric acid Phosphorous acid Phosphorous chloride Phosphorous oxychloride Phosphorous pentachloride Phosphorous pentachloride Phosphorous tribromide Phosphorous trichloride (dry) Phthalic anhydride Picric acid Potassmium aluminum sulfate Potassium bicarbonate Potassium bifluoride Potassium bisulfate Potassium bisulfite Potassium bromide Potassium carbonate Potassium chlorate Potassium chlorate Potassium chloride Potassium chromate Potassium cyanide Potassium dichromate Potassium ferricyanide Potassium fluoride Potassium hydroxide Potassium hypochlorite Potassium iodate Potassium iodide

Temp. ( C) 24 88 24–100 24–100 24–100 See Table 8 24–106 24 24 24 24 24 24–80 24 24 24 24 52 80 24 100 24 24–100 24 24–100 24–93 27 24–148 24 24 52–148 24 24 82 20 26 24 24–80 24–100 24–100 24–100 24 24 100 8 24–100 24 24–100 24–100 24–80 24–60 24 24–BPc 24

Concentration (%)b 10–100 — 10 — 10–30 — — — — — — — — — — 20–100 20–50 — 100 — 90 30 — — — 33 — — — — — 10 5.25 25 — 10–30 10 10 10–20 10–50 10–50 10 10 0.25–8.0 10–40 10–30 10–60 10–60 20 0–50 10 2–10 30

Corrosion Class B B B B B B A D C D B B B D D D A B B C B B B B B A B B A B A B B C A D B B C B C B D B B D B B B B B B D


TABLE 57.12.

803

(Continued )

Chemical Potassium metabisulfite Potassium nitrate Potassium permanganate Potassium peroxide Potassium persulfate Potassium sulfate Potassium sulfite Propionic acid Pyridine Pyridine sulfate Pyridine sulfonic acid Pyrogallic acid Quinine Quinine bisulfate Quinine tartrate Quinizarin Quinoline Quinone Saccharin solutions Salicylic acid Selenious acid þ H2SO4 þ HNO3 Silver nitrate Sodium acetate Sodium acid fluoride Sodium aluminate Sodium bicarbonate Sodium bifluoride Sodium bisulfate Sodium bisulfite Sodium carbonate Sodium carbonate Sodium chloride Sodium chlorite Sodium chromate Sodium cyanide Sodium hydrogen fluoride Sodium hydrosulfite Sodium hydroxide Sodium hypochlorite Sodium hyposulfite Sodium nitrate Sodium nitrite Sodium perborate Sodium percarbonate Sodium peroxide Sodium persulfate Sodium phosphate Sodium phosphate (tri basic) Sodium silicate Sodium sulfate Sodium sulfide Sodium sulfite Sodium tartrate Stannic chloride

Temp. ( C) 80 8 24 24 24 24–100 24 24 24–100 24 24 24 24–100 24–100 24–100 24–100 24–100 24–100 24–100 24–100 93 24 25 24 24 24 24 24–100 24–100 24 52 25 24 24–100 24 71 24 26 24 24 24 24–100 24 24 24 24 24–100 24 24 24 24–100 24–100 24 24

Concentration (%)b 10–30 0.5–10 10–40 10 10 10–20 10 10–70 10 10 20 — — 10 — — — 10 — — — 10–60 4 10 10 10 — 10–30 10 10 20 0.5–24 10 10 10 8 10–20 0–30 1 10 10 10–60 10 — 10 10 10–100 10–20 — 2–20 10–30 10–30 — 20

Corrosion Class B B C D D B B D B B A B B B B B B B B B A D B B D B B B B B D A B B B B A B C B D B D D D B B D B A A B D D (continued )

804


TABLE 57.12. (Continued ) Chemical Stannic tetrachloride (dry) Stannous bisulfate Stannous chloride Succinic acid Sulfamic acid Sulfur dioxide Sulfur trioxide Sulfuric acid Sulfurous acid Sulfuryl chloride Tanning mixtures Tannic acid Tartaric acid Tetraphosphoric acid Thionyl chloride Thiophosphoryl chloride Tetrachlorethane Titanium sulfate Titanium tetrachloride Toluene Toluene-sulfochloride Trichloroethylene Trichloronitromethane Triemanolamine Triphenyl phosphite Turpentine Vinyl chloride Zinc carbonate Zinc fluosilicate Zinc hydrosulfite Zinc sulfate Zinc chloride

Temp. ( C) 24 24–100 24 24–100 22 24–204 24 See Figure 57.4 60 24 21 24 24 24 24–149 24 63 24–100 24 24–100 24 28 24 60 28 24 24 24 21 24 35 63

Concentration (%)b

Corrosion Class

100 10 10–50 10–50 3–20 90 90

B B D B B B B — — —

20–100 30–70 10–100 — — — 10–30 — — — — — 0.4 — — 10 — 30–36 — — 25

A B B D B D B B A B B A A B C B A B D B D B B B

a

See [1]. The absence of concentration data is indicated by a dash. c Boiling point ¼ BP. b

D. GALVANIC COUPLING In acid solutions, iron is anodic to lead; therefore, the corrosion of iron is accelerated when coupled to lead. In alkaline solutions, the reverse situation applies and the corrosion of lead is accelerated by coupling with iron. Copper is anodic to lead in strong acid solutions but cathodic in alkaline solutions. In the handling of sulfuric acid and sulfates in the chemical industries, it is common practice to use Pb–Sb alloys with up to 8% Sb in pumps and valves in electrical contact with sheets and pipes of Te–Pb alloys and chemical lead. Galvanic action is inconsequential because of the formation of an insoluble insulating film of lead sulfate.

Serious galvanic corrosion does not occur under seawater with the use of caulking lead in cast iron pipe joined by bell and spigot types of joints.

E. MECHANICAL AND METALLURGICAL FACTORS The low melting point of lead allows it to recover and recrystallize at room temperature; therefore, internal stresses or work hardening may dissipate with time at ordinary temperatures. In the corrosion environments for which lead is suited, such factors are usually of minor importance, however, sustained load can result in continuous slow

805

ATMOSPHERIC CORROSION

deformation of lead, continually exposing fresh surfaces. In this case, the corrosion rate can increase. High purity lead (99.99%) may be subject to grain growth at 25 C and all alloys tend to exhibit grain growth at higher temperatures. One or more of the following elements are added to lead alloys as stabilizers or grain growth inhibitors: Cu 0.04–0.08%, Te 0.035–0.05%, Ca 0.02–0.04%. Grain growth becomes an issue if intergranular corrosion is expected. This form of corrosion can cause a significant loss in strength. F. ATMOSPHERIC CORROSION Lead has been found to be consistently durable in industrial, rural, and marine environments. The type of corrosion exhibited in each of these three environments is distinct. In rural environments, generally free of pollutants, the humidity, rainfall, and airflow are the only factors influencing corrosion

TABLE 57.13.

rate. In marine environments, chlorides are entrained in air and can exert a strong effect on corrosion. In industrial environments, sulfur oxide gases and the minerals in solid emissions can influence corrosion behavior. Despite this, lead generally gives satisfactory resistance to smoke-laden air in industrial environments and in ducts for sewer gases. Corrosion data for both lead and lead-coated steel are given in this section. F1. Lead and Lead Alloy Sheet In moist air, a dull oxide film or patina forms on the surface of lead. Oxidation generally ceases after 7 days. Several test programs have generated quantitative corrosion data in several natural atmospheres as shown in Table 57.13. The corrosion rates usually decrease with increasing exposure time, and a conservative estimate would use a linear corrosion rate over time. One example of the relationship between corrosion rate and time is shown in Figure 57.6.

Corrosion of Lead in Various Natural Outdoor Atmospheresa Corrosion Rate

Location

Type of Atmosphere

Material

Altoona, Pennsylvania Altoona, Pennsylvania New York New York Sandy Hook, New Jersey Sandy Hook, New Jersey Key West, Florida Key West, Florida La Jolla, California La Jolla, California State College, Pennsylvania State College, Pennsylvania Phoenix, Arizona Phoenix, Arizona Kure Beach, North Carolina 80-ft site Kure Beach, North Carolina 80-ft site Newark, New Jersey Newark, New Jersey Point Reyes, California Point Reyes, California State College, Pennsylvania State College, Pennsylvania Birmingham, England Birmingham, England Wakefield, England Southport, England Bourneville, England Cardington, England Cristobal CZ Miraflores CZ

Industrial Industrial Industrial Industrial Seacoast Seacoast Seacoast Seacoast Seacoast Seacoast Rural Rural Semiarid Semiarid East Coast marine East Coast marine Industrial Industrial West Coast marine West Coast marine Rural Rural Urban Urban Industrial Marine Suburban Rural Tropical marine Tropical inland

Chem Pb 1% Sb–Pb Chem Pb 1% Sb–Pb Chem Pb 1% Sb–Pb Chem Pb 1% Sb–Pb Chem Pb 1% Sb–Pb Chem Pb 1% Sb–Pb Chem Pb 1% Sb–Pb Chem Pb 6% Sb–Pb Chem Pb 6% Sb–Pb Chem Pb 6% Sb–Pb Chem Pb 6% Sb–Pb 99.96% Pb 1.6% Sb–Pb 99.995% Pb 99.995% Pb 99.995% Pb 99.995% Pb Chem Pb Chem Pb

a

See [1].

Duration (years)

mdd

mpy

10 10 20 20 20 20 10 10 20 20 20 20 20 20 2 2 2 2 2 2 2 2 7 7 1 1 1 1 8 8

0.23 0.18 0.12 0.10 0.17 0.16 0.18 0.17 0.16 0.18 0.10 0.11 0.03 0.09 0.41 0.32 0.46 0.33 0.28 0.20 0.43 0.31 0.29 0.03 0.58 0.55 0.61 0.44 0.42 0.24

0.029 0.023 0.015 0.013 0.021 0.020 0.023 0.022 0.021 0.023 0.013 0.014 0.004 0.012 0.052 0.041 0.058 0.042 0.036 0.026 0.055 0.039 0.037 0.004 0.074 0.070 0.077 0.056 0.053 0.030

806


Corrosion resistance can be further enhanced by painting, and this is widely used on terne-coated products. F3. Installation of Lead Roofing

FIGURE 57.6. Relation between corrosion and exposure time for lead exposed to marine atmosphere in Panama Canal zone. (Adapted from [2].)

The Pb–Sb alloys exhibit about the same corrosion rate in atmospheres as chemical lead; however, the greater hardness and strength of this alloy make it more desirable for use in construction applications, especially because of its reduced tendency to buckle over time. F2. Lead Coatings Effects of coating porosity on the different reactions that occur at the steel–lead interface can result in variability of corrosion behavior of lead coatings on steel. The corrosion rates of electrodeposited lead coatings in various atmospheres are given in Figure 57.7. Corrosion rates in industrial environments are much lower than the other environments. This is the reverse of experience with other metals. Depositing lead on either nickel or copper-plated steel can enhance corrosion protection. Many terne (Pb–Sn alloy) coated fuel tanks use a nickel precoat for this purpose.

Design and installation recommendations for lead roofing, flashing, and waterproofing are available in the literature from Lead Industries Association, New York, or LDA International, London. These recommendations provide for accommodation of mechanical factors and free movement of metal when thermal contraction and expansion are anticipated. Patination oil is also available that gives an immediate uniform color to exposed lead, avoiding temporary discoloration that can occur while the lead surface naturally weathers. It is a soybean alkyd resin to which solvents, a special flatting agent and driers have been added. The oil covers 60 m2/L and will dry to an even matte finish within 1 h.

G. CORROSION OF BURIED MEDIUM-VOLTAGE POWER CABLE Medium-voltage (10–50 kV) power cable is customarily placed in underground service, either directly buried in the ground or placed in buried ducts. Although almost all new installations utilize lead-free constructions, there are thousands of kilometers of lead-sheathed cable installed in many cities worldwide that are now at least 20–70 years old. The reliability of these cables, which is in many cases related to the corrosion performance of the lead cable sheathing, is of increasing concern. Several types of metallurgical or mechanical degradation mechanisms found on failed cables, including grain growth

FIGURE 57.7. Thickness loss of lead coating versus outdoor exposure. (Adapted from [2].)

LEAD ALLOYS

and fatigue, can appear to be corrosion damage and in many cases combine with corrosion activity to cause cable failure. In many cases, the only plausible replacement of the failed lead-sheathed cable is with a similar cable, and therefore it is essential to determine the cause of failure. The principal causes of cable failure corrosion are stray-current corrosion and galvanic corrosion. The causes of stray-current corrosion are manmade currents picked up by buried structures and can originate from electrically powered equipment, other power transmission lines, and nearby cathodic protection (CP) circuits. In cities, they include street railway systems, where used, welding equipment, and electroplating shops. The corrosion is localized at the point of discharge of the stray current, wherever the resistance is least for return to the power source. Pits or joined-together pits are observed in the corroded area, but the corroded lead surface is extremely smooth with almost no corrosion products [3]. In the case of straycurrent corrosion from a CP circuit, the potential of the sheath is more negative at the point of current pickup and receives the CP current. When the cable passes near the CP source, the current leaves the sheath to return to the negative and the sheath becomes anodic; this is then the site of corrosion. The products of stray-current corrosion, including oxides, chlorides, and sulfates of lead, can be carried away from the immediate reaction site by current flow, resulting in the clean corroded surfaces that are often observed. Stray-current corrosion can be prevented by reducing the electrical resistance of the offending structure or system, including the maintenance of good connections. Stray currents can also be passed from pipelines or other structures operating near high-voltage cables that cause inductive coupling; the currents produced can then be passed to nearly medium-voltage cables [3]. Galvanic corrosion of lead-sheathed cable can occur when it is grounded to a dissimilar metal structure, generally of considerable length. One example is near power stations that use large amounts of copper for grounding. Another is in connection with rusty water pipes; lead is normally cathodic to clean steel but anodic to the oxides comprising rust. Cathodic corrosion is encountered with lead-sheathed cable buried in soils containing alkali salts. In this case, lead is not removed directly by electric current but dissolved by the secondary action of the alkali produced by the current. Hydrogen ions are attracted to the metal, lose their charge, and are liberated as hydrogen gas. This results in a decrease in the hydrogen ion concentration, and the solution becomes alkaline [4]. This can be observed when a power cable passes between a structure protected by CP and the ground bed of the CP system. Lead becomes very cathodic in close proximity to the DC return, producing a final corrosion product usually consisting of bright orange lead PbO and lead–sodium carbonate. Because of the nature of their installation, differential aeration corrosion is also common on cable sheaths, but the rate of corrosion is generally rather slow because of the

807

strong tendency of the protective film formation described earlier.

H. PROTECTION MEASURES The corrosion rate of lead can be reduced by factors that help create or strengthen its protective film. Therefore, the life of lead-protected equipment can be extended, for example, by washing it with film-forming aqueous solutions containing sulfates, carbonates, or silicates. This procedure is suggested for protecting lead when it will be in contact with corrosives that do not form protective films. In water systems, the presence of organic acids whose lead salts are soluble can promote corrosion. In this case, film-forming lime or sodium silicate can be added to the water to lower corrosion rate. A pH 8–9 can be applied. This treatment is also useful to reduce release of lead from solder in copper pipe. When Portland cement is cast over a lead surface, the free alkali can cause a slow corrosive attack. This can be avoided by use of a suitable underlay such as tar, asphalt, bituminous paint, or a waterproof membrane. After 1 year the free lime in the concrete is usually sufficiently carbonated to eliminate problems; however, continued seepage of water may be a source of long-time corrosion. In chemical process equipment, the use of acid brick linings will prevent erosion effects and can lower the lead temperature, hereby reducing buckling. Use of automatic steam pressure regulation is recommended for lead coils. This allows for gradual rather than sudden introduction of steam. Similarly, avoiding use of quick shutting valves will prevent failure from water hammering. When appropriate, galvanic contact between lead and other metals should be prevented by electrical insulation between the metals. I. LEAD ALLOYS I1. Lead–Antimony (Pb–Sb) Antimony is added at levels between 1 and 13% to give greater strength, hardness, and resistance to fatigue. Cast alloys are generally harder man rolled alloys. For sheet and pipe a 6% Sb alloy is frequently employed for its higher strength. Other applications of lead alloys are shown in Table 57.14. The Pb–Sb alloys are usually not used > 95–120 C. Overaging, resulting in loss of properties, can occur at prolonged high temperatures. At compositions > 13% Sb, the alloys become too brittle for engineering uses. The tensile strength of Pb–Sb alloys varies from 17.2 MPa (2500 psi) for pure lead to 50 MPa (7280 psi) for the 13% Sb alloy. At the same time, elongation decreases from 45 to 10%, while hardness increases from 4 to 15.2 BHN (Brinell hardness number).

808


TABLE 57.14. Use of Pb–Sb Alloysa Sb (wt %)

Use

1 2–4 6 9 6–8 15 ( þ 5 Sn) 15–17 ( þ 1 Sn, 1 As)

Electric cable sheathing Storage battery connectors and parts Roofing and chemical industry Storage battery grids Lead pumps, valves, coatings Type metal, bearing metal Bearing metal

a

See [5].

I2. Lead–Tin (Pb–Sn) The most popular uses of Pb–Sn alloys are solder and terne plate. Solders can contain between 50 and 95% Sn while terne coatings for steel sheet contain between 3 and 15% Sn. The Pb–Sn alloys have much lower densities and melting points man pure Pb, however, the alloys primarily used for corrosion resistance are low in Sn such as the terne composition. Anodes of 7% Sn are used in chromium plating operations. Ternary Pb–Sn–Sb alloy can have very low coefficients of friction and have high strength. Such alloys are used in special situations such as steam spargers in chemical reactors.

resistant to corrosion in some solutions such as low-strength sulfuric acid. I5. Lead–Silver (Pb–Ag) The addition of 1% Ag to Pb gives an insoluble anode that exhibits very good corrosion resistance over a wide range of current densities. It is used for production of electrolytic zinc from strong sulfate solutions. Additions of 1% As further increase corrosion resistance, allowing use as an insoluble anode for electrowinning of manganese. The Pb–l% Ag alloy has also given good service in the cathodic protection of ships. The Pb–2.3% Ag alloy has been used as a soft solder; however, unsheltered atmospheric corrosion results in poor performance. Alloys with < 1% Ag (0.002–0.2%) have a higher corrosion resistance than pure lead in certain environments. I6. Lead–Arsenic (Pb–As) Arsenic is generally added to Pb–Sb alloys to accelerate age hardening. It also increases resistance to bending and creep of electric power cable sheathing alloys that are exposed to vibration. The most widely used alloy of this type is Pb–0.15% As–0.1% Sn–0.1% Bi.

I3. Lead–Calcium (Pb–Ca)

I7. Lead–Copper (Pb–Cu) and Other Alloys

Additions of small amounts of Ca (0.03–0.1% and Sn up to 1.5%) creates alloys with significantly improved mechanical properties and that age harden at room temperature. Moreover these alloys exhibit superior corrosion resistance to Pb–Sb alloys in many applications. Mechanical properties depend on both composition and processing. For example, a 6.4-mm (0.25-in.) section of an air-cooled static cast 0.07% Ca fully aged within 60 days of casting had a breaking strength of 37.9 MPa (5500 psi). The addition of 0.06% Sn increased strength to 51.7 MPa (7500 psi). A 1% Sn level resulted in a strength of 58.6 MPa (8500 psi). Wrought Pb–Ca–Sn alloys have improved properties with breaking strengths between 27.5 and 75.8 MPa (4000 and 11,000 psi). However, a fine grain structure is required, necessitating proper processing. The high corrosion resistance of these alloys is related to their fine grain size. The Pb–Ca alloys with or without Sn are used for battery grids, anodes, and roofing materials.

The Cu additions to Pb are limited to the ranges shown in commercial lead, up to 0.08%. Over this range Cu does not increase the corrosion resistance of lead. Additions of Ag increase corrosion resistance as noted above. Other alloying elements that have been tested for corrosion resistance in lead include Ba, Ca, Co, Au, Li, Mg, Ni, Pa, Pt, Na, Tl, and Zn. Anodes made of Ni and Tl show good performance in sulfuric acid solutions and have been adopted for electrolytic refining in some cases. Ni can be added to lead in small quantities to improve wettability of the steel during lead coating.

I4. Lead–Tellurium (Pb–Te) Additions of 0.04–0.05% Te to chemical lead containing 0.04–0.08% Cu increases corrosion resistance in several environments. These alloys have a refined grain size and exhibit work hardenability, making them useful for steam heating coil applications. The Cu-free alloys have an optimum Te composition of above 0.1%. Such alloys are less

REFERENCES 1. Lead for Corrosion Resistant Applications—A Guide, Lead Industries Association, New York, 1974. 2. F. W. Finkand W. K. Boyd,“The Corrosion of Metals in Marine Environments,” DMIC Report 245, Defense Metals Information Center, Battelle Institute, Columbus, OH. 3. S. A. Bradford, Practical Handbook of Corrosion Control in Soil, CASTI Handbook Series, CASTI Publishing, Edmonton, Alberta, 1970. 4. R. W. Hymes, “Lead Sheath Cable Corrosion—Cause and Prevention,” in Proceedings of the Fortieth Appalachian Underground Short Course, EPRI, Palo Alto, CA, 1995. 5. H. H. Uhlig (Ed.), Corrosion Handbook, Wiley, New York, 1948.

58 MAGNESIUM AND MAGNESIUM ALLOYS E. GHALI Department of Mining, Metallurgy and Materials Engineering, Laval University, Quebec, Canada

A. Production and fabrication of magnesium A1. Physical, chemical, and electrochemical properties A2. Magnesium alloys A3. Alloy designations and tempers A3.1. Cast alloys A3.2. Wrought alloys B. Corrosion forms and types B1. General corrosion and passivation B1.1. Corrosion prevention B2. Galvanic corrosion B2.1. Cathodic corrosion of aluminum B2.2. Cathodic damage to coatings B2.3. Prevention of corrosion and protection against galvanic corrosion B3. Localized attack B3.1. Pitting B3.2. Crevice corrosion B3.3. Filiform corrosion B3.4. Granular corrosion B3.5. Stress corrosion cracking B4. Corrosion fatigue C. Corrosion characteristics of magnesium alloy systems C1. Effects of alloying elements C1.1. Aluminum and the b phase C1.2. Effects of Zn and Si additions C1.3. Tramp element tolerance levels C1.4. Effect of iron and the Fe/Mn ratio C1.5. Summary C2. Influence of heat treatment C2.1. Grain refinement D. Rapid solidification D1. Effect of aluminum D2. Additional alloying elements D3. Rare earth effects

E. Magnesium finishing E1. Cleaning and surface preparation E2. Chemical and electrochemical finishing treatments E2.1. Chrome pickle and chrome-free phosphate treatments E2.2. Anodic treatments E2.3. Cathodic treatments E3. Organic finishing References

A. PRODUCTION AND FABRICATION OF MAGNESIUM The perception of magnesium as a rapidly corroding material has been a major obstacle to its growth in structural applications despite its other obviously desirable physical properties. In fact, under normal environmental conditions, the corrosion resistance of magnesium alloys is comparable or better than that of mild steel. It has been the uneducated use of magnesium in wet, salt-laden environments that has given rise to its poor corrosion reputation. Corrosion due to poor design, flux inclusions, surface contamination, galvanic couples, and incorrectly applied or inadequate surface protection schemes are all avoidable and applicable not only to magnesium but to many other metals as well. Designers and engineers in the magnesium industry have established the correct use of magnesium in corrosive environments and, over the past 30 years, have developed


809

810

MAGNESIUM AND MAGNESIUM ALLOYS

methods to improve the corrosion resistance of magnesium alloys by modifying alloy chemistry and improving surface protection technologies [1].

TABLE 58.1. Solubility Data for Binary Magnesium Alloysa Solid Solubility Element

A1. Physical, Chemical, and Electrochemical Properties Magnesium is silvery white in approach. It is a divalent metal. The atomic mass is 24.32 and the specific gravity of the pure metal is 1.738 at 20 C (68 F). The structure is closed-packed hexagonal. The lattice structure of magnesium has c/a ¼ 1.624 and atomic diameter (0.320 nm) is such that it enjoys favorable size factors with a diverse range of solute elements. The melting point is 649.5 C and the boiling point is 1107 C. The specific heat at 20 C is 1.030 kJ/kg C and the thermal conductivity at 20 C is f157.5 W/m C. The electrochemical equivalent is 0.126 mg/ C or 12.16 g/Faraday. The standard electrode potential is

EMg2 þ =Mg ¼ 2:37 V Magnesium alloys are used in the aircraft and guided weapons industries and in automotive construction because of their light weight and high strength/weight ratio. New applications are emerging because of required properties, such as high stiffness/weight ratio, ease of machinability, high damping capacity, and casting qualities. Magnesium is used as a canning material for uranium in gas-cooled reactors. Magnesium and its alloys can be used as sacrificial anodes for cathodic protection. Magnesium is itself used for alloying with other metals for different applications. The corrosion resistance of magnesium and its alloys is dependent on film formation in the medium to which they are exposed. The rate of formation, dissolution, or chemical change of the film varies with the medium and also with the metallic alloying agents, which are impurities present in the magnesium [2].

A2. Magnesium Alloys Cast magnesium alloys have always predominated over wrought alloys, particularly in Europe, where, traditionally, cast alloys have comprised 85–90% of all magnesium products. The earliest commercially used alloying elements were aluminum, zinc, and manganese, and the Mg–Al–Zn system remains the most widely used for castings. Aluminum, zinc, cerium, yttrium, silver, thorium, and zirconium are examples of widely differing metals that may be present in commercial magnesium alloys. Apart from magnesium and cadmium, which form a continuous series of solid solutions, the magnesium-rich sections of binary phase diagrams show peritectic or, more commonly, eutectic systems. Solubility data for binary magnesium alloys are given in

Lithium Aluminum Silver Yttrium Zinc Neodymium Zirconium Manganese Thorium Cerium Cadmium Indium Thallium Scandium Lead Thulium Terbium Tin Gallium Ytterbium Bismuth Calcium Samarium Gold Titanium

at %

wt %

17.00 11.80 3.80 3.75 2.40 1.00 1.00 1.00 0.52 0.10 100.00 19.40 15.40 15.00 7.75 6.30 4.60 3.35 3.10 1.20 1.10 0.82 1.00 0.10 0.10

5.50 12.70 15.00 12.50 6.20 3.00 3.80 2.20 4.75 0.50 100.00 53.20 60.50 24.50 41.90 31.80 24.00 14.50 8.40 8.00 8.90 1.35 6.40 0.80 0.20

a

Reproduced with permission [3, 4], ASM International.

b

Solid solubility ¼ SS.

Systemb Eutectic Eutectic Eutectic Eutectic Eutectic Eutectic Peritectic Peritectic Eutectic Eutectic Complete SS Peritectic Eutectic Peritectic Eutectic Eutectic Eutectic Eutectic Eutectic Eutectic Eutectic Eutectic Eutectic Eutectic Peritectic

Table 58.1; the first 10 elements are those used in commercially available alloys. Although early Mg–Al–Zn castings suffered severe corrosion in wet or moist conditions, the corrosion performance was significantly improved as a result of the discovery, in 1925, that small additions (0.2%) of manganese gave increased resistance. With this element, iron and certain other heavy metal impurities formed relatively harmless intermetallic compounds, some of which separate out during melting. In this regard, the classic work by Hanawalt et al. [5] showed that the corrosion rate increased abruptly once tolerance limits were exceeded; these tolerance limits are 5, 170, and 1300 ppm for nickel, irons, and copper, respectively. The corrosion rate of pure magnesium as a function of iron content is shown in Figure 58.1, which clearly illustrates the tolerance limit for iron. Another problem with earlier magnesium alloy castings was that grain size tended to be large and variable, often resulting in poor mechanical properties, microporosity, and, in wrought products, excessive directionality of properties. Values of proof stress also tended to be low relative to tensile strength.

PRODUCTION AND FABRICATION OF MAGNESIUM

FIGURE 58.1. Effect of iron on corrosion of pure magnesium; alternate immersion test in 3% NaCl [5].

In 1937, it was discovered that zirconium had an intense grain refining effect on magnesium. The lattice parameters of zirconium are very close to those of magnesium. Paradoxically, zirconium could not be used in most existing alloys at that time because it was removed from solid solution owing to the formation of stable compounds with aluminum and manganese. This problem led to the evolution of a completely new series of cast and wrought zirconium-containing alloys with much improved mechanical properties at both room and elevated temperatures. Alloys containing zirconium as a grain-refining agent have the iron content reduced to 0.004% because impurities separate during the alloying procedure. These alloys are now widely used in the aerospace industries. A3. Alloy Designations and Tempers A3.1. Cast Alloys. Two major cast alloy magnesium systems are available to the designer. The first system includes alloys containing 2–10% Al, combined with minor additions of zinc and manganese. These alloys are widely available at moderate cost, and their mechanical properties are satisfactory from 95 to 120 C (200 to 250 F). At higher temperatures, the properties deteriorate. The second system consists of magnesium alloyed with various elements (rare earths, zinc, thorium, silver, etc.) except aluminum, all containing a small but important zirconium content that imparts a fine grain structure that improves mechanical properties. These alloys generally possess much better elevated temperature properties, but their more costly elemental additions, combined with the specialized manufacturing technology required, result in significantly higher costs. Many of the casting alloys are given simple heat treatments to improve

811

their properties, while the wrought alloys can be obtained in a number of tempers [6]. An international code for designating magnesium alloys does not exist, although there has been a tendency toward adopting the method used by the American Society for Testing and Materials (ASTM) B275-94 [7–9]. In this system, the first two letters indicate the principal alloying elements according to the following code: A aluminum, B bismuth, C copper, D cadmium, E rare earths, F iron, H thorium, K zirconium, L lithium, M manganese, N nickel, P lead, Q silver, R chromium, S silicon, T tin, W yttrium, Y antimony, and Z zinc. The letter corresponding to the element present in greater quantity in the alloy is used first; if they are equal in quantity the letters are listed alphabetically. Letters are followed by numbers that represent the nominal compositions of these principal alloying elements in weight percent, rounded off to the nearest whole number, (e.g., AZ91 indicates the alloy Mg–9Al–1Zn), the actual composition ranges being 8.3–9.7% Al and 0.4–1.0% Zn. Suffix letters A, B, C are chronologically assigned and usually refer to purity improvement. The X is reserved for experimental alloys. For heat-treated or work-hardened conditions, the designations are specified by the same system as that used for aluminum alloys [8]. Commonly used tempers are T5, alloys artificially aged after casting; T6, alloys solution treated, quenched, and artificially aged; and T7, alloys solution treated and stabilized [10]. Because of the particularly high solid solubility of yttrium in magnesium (12.5% max) and the amenability of Mg–Y alloys to age hardening, a series of Mg–Y–Nd–Zr alloys has been produced, which combine high strength at ambient temperatures with good creep resistance at temperatures up to 300 C [11, 12]. The heat-treated alloys have a resistance to corrosion, which is superior to that of other high-temperature magnesium alloys and comparable to many aluminum-based casting alloys [13, 14]. Since pure yttrium is expensive and difficult to alloy with magnesium because of its high melting point (1500 C) and its strong affinity for oxygen, a cheaper yttrium-containing (75% Y) mischmetal together with heavy rare earth metals such as gadolinium and erbium could be substituted for pure yttrium (Table 58.2) [15]. A3.2. Wrought Alloys. Wrought materials are produced mainly by extrusion, rolling, and press forging at temperatures in the range 300–500 C. As with cast alloys, the wrought alloys may be divided into two groups according to whether or not they contain zirconium (Table 58.3). Specific alloys have been developed that are suitable for wrought products, most of which fall into the same categories as the casting alloy already [9]. Examples of sheet and plate alloys are AZ31 (Mg–3 Al–1 Zn–0.3 Mn), which are the most widely used because they offers a good combination of strength, ductility, and corrosion resistance, and thoriumcontaining alloys such as HM21 (Mg–2 Th–0.6 Mn), which

812

TABLE 58.2.

Nominal Composition, Typical Tensile Properties, and Characteristics of Selected Magnesium Casting Alloysa Nominal Composition, wt %

Tensile Properties

b

Rare Earth

ASTM Designation

British Designation

AZ63

Si

Cu

Zr

MM

Nd

Th

Y

Ag

Condition

0.2% Proof Stress (Mn/m2)

Tensile Strength (Mn/m2)

Elb (%)

Al

Zn

Mn

6

3

0.3

As sand cast T6

75 110

180 230

4 3

Good room temperature strength and ductility

Characteristics

AZ81

A8

8

0.5

0.3

As sand cast T4

80 80

140 220

3 5

Tough, leak tight casting with 0.0015% Be used for pressure die casting

AZ91

AZ91

9.5

0.5

0.3

As sand cast T4 T6 As chill cast T4 T6

95 80 120 100 80 120

135 230 200 170 215 215

2 4 3 2 5 2

General purpose alloy used for sand and die casting

AM50

5

0.3

As die cast

125

200

7

High pressure die castings

AM20

2

0.5

As die cast

105

135

10c

Good ductility and impact strength

AS41

4

0.3

1

As die cast

135

225

4.5c

Good creep properties up to 150 C

AS21

2

0.4

1

As die cast

110

170

4c

Good creep properties up to 150 C

ZK51

ZK61

ZSZ

4.5

0.7

T5

140

235

5

Sand castings, good room temp, strength and ductility

6

0.7

T5

175

275

5

As for ZK51

ZE41

RZ5

4.2

ZC63

ZC63

6

EZ33

ZREI

2.7

0.7

0.5

1.3

3

0.7

3.2

T5

135

180

2

Sand castings, good room temp. strength and castability

T6

145

240

5

Pressure tight castings, good elevated temp., strength, weldable

Sand cast T5

95 100

140 155

3 3

Good castability, pressure tight, weldable, creep resistant up to 250 C

Chill cast T5 HK31

MTZ

0.7

3.2

Sand cast (T6) 90

185

4

Sand casting, good castability, weldable, creep resistant up to 350 C

HZ32

ZTI

0.7

3.2

Sand or chill cast (T5)

90

185

4

As for HK31

QE22

MSR

0.7

2.5

2.5

Sand or chill cast (T6)

185

240

2

Pressure tight, weldable, high proof stress up to 250 C

QH21

QH21

0.7

1

2.5

As sand cast (T6)

185

240

2

Pressure tight, weldable, good creep resistance and proof stress up to 300 C

WE54

WE64

0.5

3.25d

5.1

T6

200

285

4c

WE54

WE43

0.5

3.25d

4

T6

190

250

7c

High strength at room and elevated temp, good corrosion resistance, weldable

a

2.2

1

Reproduced with permission from [15], Institute of Materials, London, UK. Mischmetal ¼ MM; El ¼ elongation. c Values quoted for tensile properties are for separately cast test bars and may not be realized in certain parts of castings. d Contains some heavy metal rare earth elements. b

813

814


TABLE 58.3. Nominal Composition, Typical Tensile Properties, and Characteristics of Selected Magnesium Casting Alloysa Nominal Composition, wt % ASTM Designation

British Designation

MI

AM503

AZ31

AZ31

Al Zn

3

1

Mn

200 230 200

4 4 4

0.3

Sheet, plate/O

120

240

11

0.2b

Sheet, plate/H24 Extrusions/F Forgings/F Extrusions/F Forgings/F Forgings/T6

160 130 105

250 230 200

6 4 4

180 160 200

260 275 290

7 7 6

High-strength alloy, weldable High-strength alloy

Sheet, plate/O

120

240

11

165 155 125

250 235 200

6 8 9

Medium-strength alloy, good formability, good damping capacity

300 95

325 115

3 10

Extrusions/T5 Forgings/T5 Extrusions/F Extrusions/T5 Forgings/T5 Sheet,

210 205 210 240 160 170

295 290 285 305 275 230

8 7 6 4 7 4

Plate/H24

180

255

4

Extrusions/T5 Sheet, plate/T8

135

215

6

Sheet, plate/T81 180

255

4

Forgings/T5

175

255

3

Extrusions/F

120

215

7

Forgings/F

130

230

6

AZ80

AZ80

ZM21

ZM21

8.5 0.5 0.12b 2 1

ZK31

ZW3

ZK61

0.3 0.15b 0.2

0.75 0.15b

0.6

6

0.8

HK31

HZ11

Sheet, Plate/H24 Extrusions/F Forgings/F 1.25 Extrusions/T6 14 Sheet, plate/T7

3

0.7 3.2

HM21

0.8

ZTY

0.6

El (%) Characteristics

70 130 105

6.5 1

6.5

Tensile Strength (Mn/m2)

Sheet, plate/F Extrusions/F Forgings/F

AZM

1.2

Li Condition

0.2% Proof Stress (Mn/m2)

1.5

AZ61

ZMC711 LA 141

Zr Th Cu

Tensile Properties

2

0.6 0.8

Low- to mediumstrength alloy, weldable, corrosion resistant Medium-strength alloy, weldable, good formability

High-strength alloy Ultralight weight (specific gravity 1.35) High-strength alloy, some weldability High-strength alloy

High creep resistance up to 350 C, short time High creep resistance up to 350 C, short time exposure up to 425 C weldable Creep resistance up to 350 C, weldable

a

Reproduced with permission from [15], Institute of Materials, London, UK. Minimum.

b

show good creep resistance at temperatures of up to 350 C. Magnesium alloys can be extruded at temperatures > 250 C into either solid or hollow sections at speeds that depend on alloy content. Higher strength alloys such as AZ81 (Mg– 8

Al–1 Zn–0.7 Mn), ZK 61 (Mg–6 Zn–0.7 Zr), and the more recent composition ZCM711 (Mg–6.5 Zn–1.25 Cu–0.75 Mn) all have strength/weight ratios comparable to those of the strongest wrought aluminum alloys. The alloy ZM21

CORROSION FORMS AND TYPES

(Mg–2 Zn–1Mn) can be extruded at high speeds and is the lowest cost magnesium extrusion alloy available. Again, thorium-containing alloys, such as HM31 (Mg–3 Th– l Mn), show the optimal elevated temperature properties. Magnesium forgings are less common and are almost always press formed rather than hammer forged.

B. CORROSION FORMS AND TYPES B1. General Corrosion and Passivation Magnesium exposed to air is covered by a gray oxide film, which protects the metal from further oxidation. Magnesium can be heated in air to the melting point without burning. However, the fine divided metal reacts vigorously by heating or by contact with water or humid atmospheres. The evolved hydrogen from this reaction can result in an explosive mixture. In aqueous solutions, magnesium dissociates by electrochemical reaction with water to produce a crystalline film of magnesium hydroxide, Mg(OH)2 [16], and hydrogen gas, a mechanism, which is highly insensitive to the oxygen concentration [17]. Subsequently, all that is needed for rapid corrosion are sites of easy hydrogen discharge [18]. The probable primary overall corrosion reaction for magnesium in aqueous solutions is MgðsÞ þ 2H2 Oð‘Þ Y MgðOHÞ2 ðsÞ þ H2 ðgÞ

whether it is from the properties of the film or the evolution of hydrogen gas. The Pilling/Bedworth ratio for Mg(OH)2 is 1.77, which indicates a resistant film in compression. A combination of internal stresses and easy basal cleavage may account for a portion of the cracking and curling of the film. Thus, the structure of the corrosion product directly influences the corrosion behavior of the base metal [19]. Magnesium may form a surface film, which protects it in alkaline environments, and poorly buffered environments where the surface pH can increase. Passivity of magnesium is destroyed by several anions, including chloride, sulfate, and nitrate. Alloying affects the nature of this film, but these effects are poorly understood. The corrosion of magnesium and its alloys is strongly dependent on the absence of impurity elements, some of which have well-defined tolerance levels above which corrosion resistance drops dramatically. For conventional magnesium alloys, these tolerance limits must be observed even if extensive surface treatments are applied [20]. The Pourbaix (potential–pH) diagram [21] shows possible protection of magnesium at high pH values, which may result from Mg(OH)2 formation during the corrosion reaction. Perrault [22] considered the formation of MgH2 and Mg þ and assumed that thermodynamic equilibrium cannot exist for a magnesium electrode in contact with aqueous solutions. Such equilibrium is, however, possible if the hydrogen overpotential is about 1 V and the pH is > 5. The following reactions are considered in the E–pH diagram (Fig. 58.2): 2H þ þ 2e ! H2

This overall reaction can be described in terms of anodic and cathodic reactions as follows: Anodic Reaction :

ð58:1Þ

MgH2 ! Mg2 þ þ H2 þ 2e

MgH2 þ 2OH ! MgðOHÞ2 þ H2 þ 2e

Mg Y Mg2 þ þ 2e ðdissolution of MgÞ

Mg

2þ

ð58:2Þ

þ 2OH ! MgðOHÞ2

þ

MgðsÞ þ 2ðOHÞ Cathodic Reaction :

Y MgðOHÞ2 ðsÞ þ 2e þ

2H þ 2e

Y H2 ðgÞ

Mg þ 2OH ! MgðOHÞ2 þ 1e þ

ð58:3Þ ð58:4Þ

Mg þ ! Mg2 þ þ e

and/or

815

ð58:5Þ þ

Mg þ 2H2 O ! MgðOHÞ2 þ 2H þ e MgH2 ! Mg þ þ H2 þ e

ð58:6Þ

ð58:7Þ

ðevolution of hydrogen gasÞ A subsequent reaction giving OH ions can occur and/or 2H2 O þ 2e Y H2 ðgÞ þ 2ðOHÞ In general, the magnesium corrosion products resulting from the anodic reaction depend on the environment and may include carbonate, hydroxide, sulfite, and/or sulfate compounds. The hydroxide film, brucite, has a hexagonal crystalline structure that is layered, alternating between Mg and hydroxide ions, facilitating easy basal cleavage. Cracking and curling of the film have been noted though it is not clear

Although magnesium has a standard electrode potential at 25 C of 2.37 V, its corrosion potential is more negative than 1.5 V in dilute chloride solution or a neutral solution with respect to the standard hydrogen electrode due to the polarization of the formed film of Mg(OH)2. The oxide film on magnesium offers considerable surface protection in rural and some industrial environments, and the corrosion rate of magnesium lies between that of aluminum and that of lowcarbon steels (Table 58.4). In natural atmospheres, the corrosion of magnesium can be localized. The conductivity, ionic species, temperature of the electrolyte, alloy composition and homogeneity,

816


FIGURE 58.2. Equilibria of Mg–H2O system in presence of H2 at 25 C. TABLE 58.4. Results of 2.5-Year Exposure Tests on Sheet Alloysa

Material

Corrosion Rate (mm/year)

Loss of Tensile Strength after 2–5 years (%)

2.0 18.0 150.0

2.5 7.4 75.4

2.0 27.7 25.4

1.5 11.2 11.9

0.1 13.0 15.0

0.4 5.9 7.5

Marine Atmosphere Aluminum alloy 2024 Magnesium alloy AZ31 Low-carbon steel (0.27%C) Industrial Atmosphere Aluminum alloy 2024 Magnesium alloy AZ31 Low-carbon steel (0.27%C) Rural Atmosphere Aluminum alloy 2024 Magnesium alloy AZ31 Low-carbon steel (0.27%C) a

attack from any source can jeopardize the satisfactory performance of magnesium components, attack resulting from galvanic corrosion is probably the most detrimental. Because magnesium is anodic (or sacrificial) to all other engineering metals, the severe corrosive attack that often occurs with magnesium assemblies in saltwater environments has long been a deterrent to the use of magnesium alloys in structural applications. Galvanic corrosion of magnesium alloys can generally be attributed to two basic causes: (1) poor alloy quality due to excessive levels of heavy metal or flux contamination and (2) poor design and assembly practices, which can result in severe galvanic corrosion attack. With the recent development of fluxless melt protection and new high-purity alloys, such as AZ91D, AZ91E, AM60B, and others, a renewed interest in magnesium has developed due to the improved corrosion resistance of these alloys. The new alloys offer no defense against galvanic corrosion attack; however, their improved performance in assemblies can only be realized if proper measures are taken to control the potential for galvanic attack through careful design, selection of compatible materials, and the selective use of coatings, sealants, and insulating materials. The severity of galvanic activity is determined by the galvanic current which flows in the completed circuit. This can be expressed as follows: Ea þ Re I ¼ Ek Rm where Ek and Ea are the polarized measured potentials of the cathode and anode, respectively, and Rm and Re are the resistance of the metal-to-metal contact and the electrolyte portions of the circuit, respectively. The electrochemical reactions are Anode reaction : MgðmetalÞ ! Mg2 þ þ 2e ð58:8Þ Cathode reaction : 2H2 O þ 2e ! H2 þ 2OH

ð58:9Þ


differential aeration, and so on influence the corrosion morphology. B1.1. Corrosion Prevention. Effective corrosion prevention for magnesium components and assemblies begins at the design stage. General corrosion attack in saltwater exposures can be minimized through the selection of highpurity magnesium alloys cast without introducing heavy metal contaminants and flux inclusions B2. Galvanic Corrosion For continuous outdoor use, where magnesium assemblies may be wet or subjected to salt splash or spray, precautions against galvanic attack must be taken. Although corrosive

In many practical applications, Rm is negligibly small due to mechanical, electrical, or cost requirements, and Re (the electrolyte resistance) becomes the controlling factor in the circuit resistance. If the environment is rich in marine mists or deicing salts, the use of drain holes and sealants can help control corrosion by forcing the galvanic current in the electrolyte to flow through a thin and, therefore, highly resistive film (maximizing Re). In practice, this limits the galvanic activity to an area 3.2–6.4 mm (18 -- 14 in.) wide on either side of the magnesium–cathode interface. The magnesium alloy can still suffer severely; however, the cathode does not polarize sufficiently to reduce or eliminate the effective potential difference (Ek Ea) [24]. The degree to which the corrosion of magnesium is accelerated by the galvanic couple in a given environment (i.e., a given Re) depends in part on the relative positions of


the two metals in the electrochemical series. Equally important, the polarization that reduces the potential difference of the couple as the galvanic current develops. Because magnesium shows little, if any, anodic polarization in saltwater exposures, the reduction of the potential difference in the galvanic cell typically results from polarization of the cathode, where water is reduced to hydrogen gas and hydroxyl ion [Eq. 58.9]. Some metals, such as iron, nickel, and copper, serve as efficient cathodes in what is thought to be the stepwise process of accepting and reducing hydrogen ion to an atomic form (H) where it then combines to form the evolved hydrogen gas (H2). These metals have a low hydrogen overvoltage and can consequently cause severe galvanic corrosion of magnesium. Other metals, such as aluminum, zinc, cadmium, and tin, while equally cathodic to magnesium in some environments, serve as much less effective cathodes due to their tendency to inhibit the combination of atomic hydrogen on surfaces to form the hydrogen gas that evolves. Data on galvanic corrosion of magnesium alloys were compiled in tests at Kure Beach, NC, in which sheets of dissimilar metals were fastened to panels of AZ31B and AZ61A. The dissimilar metals were divided into five groups

817

(Table 58.5), ranging from the recommended Group 1 to Group 5, metals in Group 5 caused severe galvanic corrosion of magnesium alloys. Aluminum alloys containing small percentages of copper (7000 and 2000 series and 380 die-casting alloy) may cause serious galvanic corrosion of magnesium in saline environments. Very pure aluminum is quite compatible, acting as a polarizable cathode; but when iron content exceeds 200 ppm, cathodic activity becomes significant (apparently because of the depolarizing effect of the intermetallic compound FeAl3), and galvanic attack of magnesium increases rapidly with increasing iron content. The effect of iron is diminished by the presence of magnesium in the alloy. This agrees with the relatively compatible behavior of aluminum alloys 5052, 5056, and 6061 shown in Table 58.6 [23]. The corrosion of magnesium being largely cathodically controlled, the polarization characteristics of the coupled cathode will largely control the galvanic corrosion. In a highly conducting medium, such as 3% NaCl, most metals will not polarize to the magnesium potential until a relatively high current density is reached. In contact with metals such as steel or nickel, very high corrosion currents are obtained in most

TABLE 58.5. Relative Effects of Various Metals on Galvanic Corrosion of Magnesium Alloys AZ31B and AZ61A Exposed at the 24.4- and 244-m (80- and 800-ft) Stations, Kure Beach, NCa Group 1 (Least Effect)

Group 2

Group 3

Group 4

Group 5 (Greatest Effect)

Al alloy 5052 Al alloy 5056 Al alloy 6061

Al alloy 6063 AlClad alloy 7075 Al alloy 3003 Al alloy 7075

AlClad alloy 2024 Al alloy 2017 Al alloy 2024 Zinc

Zn-plated steel Cd-plated steel

Low-carbon steel Stainless steel Monel, titanium Lead, copper Brass

a

See [18].

TABLE 58.6.

Corrosion of Mg 6% Al–3% Zn 0.2% Mn Alloy Galvanically Connected to Other Metals in Various Mediaa Corrosion Rate (mdd) 3% NaCl

Midland Tap Water

Distilled Water

0.35 cm

0.35 cm

Separation Dissimilar Metal

Close Contact

0.35 cm

2.0 cm

Steel 23,400 25,500 8300 Aluminum alloys 2024 12,800 25,700 6800 Nickel 18,800 22,400 6600 Aluminum alloys 1100 14,500 15,600 4100 Copper 8500 8200 3700 Brass 7100 4000 2500 Aluminum alloys 5056 1900 Cd-plated steel 5200 2200 1000 Zinc 6200 1300 900 Mg–l.5%Mn 50 Mg–6% Al–3% Zn–0.2% Mn 200 Size of specimen 4 1.3 0.2 cm (1.5 0.5 0.079 in.) Relative areas 1:1 (mounted face to face). Surface preparation Aloxite 150 ground Velocity Quiescent Duration of test: 3% NaCl, 3 h; midland tap water, 24 h; distilled water, 4 days. a

See [2].

10 cm 3900 3200

300 90 210 40 90 1700 60 10 40 700 30 2 7 Temperature Aeration Volume of testing solution

18 6 19 4 15 14 3 14 8 3 3 Room Nat. convect 100 mL

818


highly conducting media. An exception is the Al–5% Mg rivet alloy that normally polarizes at a very low current density. The conductivity and composition of the medium in which a couple is immersed are controlling factors in the rate of galvanic corrosion. Equal areas of various cathodic materials and a magnesium alloy were tested by continuous immersion in 3% NaCl, Midland tap water containing approximately 70 ppm chloride, and in distilled water (Table 58.6). All commonly used metals cause galvanic corrosion of magnesium in a strong chloride electrolyte. Cadmium or zinc plating of the more cathodic metals, such as iron or steel, reduces the galvanic corrosion to one tenth the rate; a reduction in the conductivity, for example, a change from 3% NaCl to tap water, causes an even greater reduction in galvanic corrosion rate. Under conditions where the corrosion product is not continuously removed or under conditions of high cathodic

current density where the surroundings may become strongly alkaline, both the magnesium and an amphoteric contacting metal such as aluminum may suffer severe attack. Aluminum alloys containing appreciable magnesium, such as 5052, 6053, 5056, are least severely attacked in chloride media when galvanically coupled. This fact was observed in galvanic couples of magnesium and aluminum alloys exposed to tide water and in the atmosphere at Hampton Roads, VA. Zinc, cadmium, or tin plating on steel all reduce galvanic attack of magnesium substantially when compared to that produced by uncoated steel. The relative compatibility of the electroplates in descending order has generally been concluded to be tin, cadmium, and zinc. This is consistent with the data presented at Figure 58.3, where the compatibility of various fasteners (plated on coated steel, plus alternative materials) was determined by the magnesium weight loss in a

FIGURE 58.3. Galvanic corrosion produced by dissimilar fasteners in AZ91D magnesium alloy. (Reproduced with permission from [24].)


10-day salt spray exposure. Certain zinc and aluminum-filled polymer coatings on steel actually produced more damage to the magnesium than bare steel. This effect may be due to either an increase in the active cathode surface area resulting from the fine metal powders or flake employed in the coatings or it may be due to the presence of a catalytic contaminant on the metal powder surface, such as iron. An inorganic chromate treatment on cadmium electroplate (and perhaps on other electroplates, and on metal surfaces) was as effective in reducing the galvanic attack on magnesium as an epoxy coating. This observation is consistent with the known inhibitive effect of chromates on the cathodic reduction process. The salt spray test is thought to produce a result biased against zinc due to the rapid cathodic attack on the zinc electroplate produced in this severe test exposure. This attack does not occur in many natural environments. The most compatible fastener coatings are based on zinc plating, with modifications to extend the life of the zinc. These modifications include chromating, silicate treatments, and alloying with tin [25]. In salt spray tests using cast iron disks coupled to AZ91 die cast plates and separated by plastic spacers, it was found that a separation of about 4.45 mm (175 mils) was needed to ensure the absence of galvanic corrosion (Fig. 58.4) [24]. The relative areas of the magnesium anode and the dissimilar metal cathode have an important effect on the galvanic corrosion damage that occurs. A large cathode coupled to a small area of magnesium results in rapid penetration of the magnesium because the galvanic current density at the small magnesium anode is very high, and anodic polarization in chloride solutions is very limited. Painted magnesium should not be coupled to an active cathodic metal if the couple will be exposed to saline or

819

aggressive environments. A small break in the coating at the junction results in a high concentration of galvanic current at that point. Unfavorable area effects can also be seen in the behavior of some proprietary coatings using aluminum or zinc powder [23]. B2.1. Cathodic Corrosion of Aluminum. Aluminum can be attacked by the strong alkali generated at the cathode when magnesium corrodes sacrificially in static NaCl solutions. Such attack destroys compatibility in alloys containing significant iron contamination, apparently by exposing fresh, cathodic active sites with low overvoltage. The aluminum alloys, having substantial magnesium content (5052 and 5056) are more resistant to this effect but not completely so. A 5052 alloy would meet the essential requirement for a fully compatible aluminum alloy with a maximum of 200 ppm Fe or a 5056 alloy with a maximum of 1000 ppm Fe [23]. Cathodic corrosion of aluminum is much less severe in seawater than in NaCl solution because the buffering effect of magnesium ions reduces the equilibrium pH from 10.5 to 8.8. The compatibility of aluminum with magnesium is, accordingly, better in seawater and is less sensitive to iron content [26]. Aluminum oxide is amphoteric, that is, soluble in alkaline as well as acid solution. The standard potentials of these two half-reactions are Al3 þ þ 3e ¼ Al ð 1:66 VÞ

Acid Alkaline

H2 AlO3

þ H2 O þ 3e

ð58:10Þ

¼ Al þ 4OH ð2:35 VÞ ð58:11Þ

Half-reaction (58.11) has nearly the same standard potential as that for acidic dissolution of magnesium: Mg2 þ þ 2e ¼ Mg ð 2:37 VÞ Commercial aluminum alloys contain several thousand parts per million (ppm) of iron in the form of the intermetallic FeAl3. The mutually destructive galvanic action between magnesium and commercial aluminum alloys in salt water proceeds as follows:

FIGURE 58.4. Effect of spacer thickness on the galvanic corrosion of AZ91 magnesium coupled to cast iron disks through plastic spacers. (Reproduced with permission from [24].)

1. Rise in the pH of the liquid in contact with the aluminum member. This is most likely the result of galvanic current flow between the magnesium and the initially passive aluminum. 2. Shift of the aluminum potential in the active direction in accordance with the half-reaction (58.11) 3. Exposure of iron aluminum intermetallic particles (e.g., FeAl3, which then engage in separate galvanic activity with the magnesium). This galvanic current flow accounts for the severe sacrificial corrosion of the

820


magnesium, and the alkali generated at the cathode ensures continued corrosion of the aluminum in accordance with half-reaction (58.11) [25]. B2.2. Cathodic Damage to Coatings. Hydrogen evolution and strong alkalinity generated at the cathode can damage or destroy organic coatings applied to fasteners or other accessories coupled to magnesium. Alkali-resistant resins are necessary, but under severe conditions, such as salt spray or salt immersion, which do not simulate adequately a real application, the coatings may be simply blown off by hydrogen, starting at small voids or pores. B2.3. Prevention of Corrosion and Protection against Galvanic Corrosion [24]. Use Indoors and in Sheltered Outdoor Environments. For indoor use, where condensation is not likely, no protection is necessary. Even in some sheltered outdoor environments, unprotected magnesium components can give good service life providing the absence of water traps, good ventilation, warm component temperature, or the presence of an oil film, and so on. Seal Faying Surfaces. Sealing compounds, such as nonacidic silicone RTVs, polysulfide, epoxy resins, or plastic tapes, can be employed. If possible the compound or tape should extent beyond the joint interface by 3.2–9.5 mm (18 -- 38 in.) (Fig. 58.5). Iron, and to a lesser extent zinc, based phosphate treatments can replace chromates as inhibitors since the discharge or the formation of CR(VI) in natural water should be avoided. The joining of two magnesium components invariably involves the use of dissimilar metal fasteners and the formation of a crevice at the joint. Good engineering practice dictates that for corrosive conditions some precautions must be taken (Fig. 58.5). Magnesium faying or mating surfaces should be assembled using “wet assembly” techniques. Inhibited primers or

FIGURE 58.5. Schematic of a method to protect faying surfaces in magnesium-to-magnesium assemblies. (Reproduced with permission from [24].)

sealing compounds are placed between the surfaces at the time of assembly. Sealing/jointing compounds of the polymerizing or nonpolymerizing type are preferred as they will remain flexible and resist cracking. Polymerizing-type compounds are also used for caulking operations. In bolted assemblies, the retorquing of bolts a short while after assembly helps to eliminate any joint relaxation problems. For additional protection, mating surfaces can be primed prior to assembly and overpainted after assembly. Joining Magnesium to Dissimilar Metal Assemblies. Good design can play a vital role in reducing galvanic corrosion (Fig. 58.6). The elimination of a common electrolyte may be possible by the provision of a simple drain or shield to prevent liquid entrapment at the dissimilar metal junction (Fig. 58.7). Alternatively, the location of screws or bolts on raised bosses may also help avoid common electrolyte contact, as would the use of nylon washers, spacers, or similar moistureimpermeable gaskets. The use of studs in place of bolts will, provided the captive ends of the studs are located in blind holes, reduce the area of dissimilar metal exposed by up to 50%. The use of wet assembly techniques will eliminate galvanic corrosion crevices. Caulking the metal junctions will, by lengthening the electrolytic path, increase the electrical resistance (Re) of the galvanic couple and so reduce the degree of attack should it occur (Fig. 58.7). Vinyl tapes have also been used to separate magnesium from dissimilar metals or a common electrolyte and so prevent galvanic attack (Fig. 58.8). Finally, overpainting the magnesium and more importantly the dissimilar metal after assembly will effectively insulate the two materials externally from any common electrolyte. Use Compatible Materials. Contacting components, fasteners and inserts, and so on should be chosen for their compatibility; for example, a nonconductive, nonporous material; 5000 or 6000 series aluminum alloys; or Sn , Cd , or Zn-plated ferrous alloys. The compatibility of plated fasteners can be further improved by the use of aluminum washers, organic coatings, or other inhibiting films. Dissimilar metals that are compatible with magnesium are the aluminum–magnesium (5000 series) or aluminum–magnesium–silicon (6000 series) alloys, which should be used for washers, shims, fasteners (rivets and special bolts), and structural members where possible. Aluminum, zinc, cadmium, and tin are used to coat steel or brass components in order to reduce the galvanic couple with magnesium, under mild corrosive environments, but will have minimal effect in corrosive conditions, where additional precautions are required.


821

FIGURE 58.6. Design consideration for reducing galvanic corrosion: (a) proper versus improper bolt location, (b) poor sealing practice, (c) good sealing practice, and (d) good sealing practice when direct metal-to-metal contact is required. (Reproduced with permission from [24].)

FIGURE 58.7. Examples of good practice for bushing installations. (Reproduced with permission from [24].)

FIGURE 58.8. Proper use of insulating tapes to avoid galvanic corrosion (Reproduced with permission from [24].)

822


If painting is to be employed on only one of the contacting components, paint the cathodic material. Painting both components is a better practice. Paints employed on cathodic components and complete magnesium assemblies should be chosen for resistance to alkalis in order to prevent stripping of the coating. Joining Magnesium to Nonmetallic Assemblies. Joining magnesium-to-wood assemblies presents an unusual problem because of the water absorbency of wood and their tendency to leach out natural acids. To protect magnesium from attack, the wood should first be sealed with paint or varnish and the faying surface of the magnesium should be treated as magnesium-to-magnesium assemblies. The joining of magnesium to carbon fiber reinforced plastics, in the presence of a common electrolyte, could result in corrosion of the magnesium unless similar assembly precautions were observed. High-Purity Alloys. Under corrosive conditions, the use of high-purity magnesium alloys will not reduce the effects of galvanic corrosion significantly. B3. Localized Attack General corrosion can lead to localized corrosion, which is favored by a weak electrolyte and small anode/cathode relative area ratios. Localized attack takes the form of pitting, crevice, and filiform corrosion. Intergranular corrosion can also be considered as a localized attack due to metallurgical structure; however, in the case of magnesium alloys, this type of corrosion can be named more properly as granular attack. B3.1. Pitting. When corrosion occurs on a smooth machined magnesium alloy surface, this surface is roughened by the chemical action, and after the initial attack the degree of roughness does not change appreciably. In atmospheric attack the roughening is really a microscopic form of pitting. There is a noticeable difference between the appearance of the aluminum-containing magnesium-rich alloys and the zinc/zirconium-containing magnesium alloys. In the former, the microscopic pits in the surface exposed to the weather tend to be narrow and relatively deep, whereas in the latter they are wider and tend to overlap, leading to a slightly wavy appearance [6]. In the usual industrial atmospheric conditions the attack is uniform, but in immersed conditions, including corrosion under pools of condensate, attack may be, and usually is, irregular; some areas become anodic to other areas and, as corrosion proceeds at the anodic areas, a pitting develops. The unequal attack, which occurs in tap water, condensate, and other mild electrolytes, may lead to perforations of thingauge sheet and even to deep pitting of castings. In stronger electrolytes, the effect is variable. In chloride solutions, such as seawater, attack on the metal usually results in pitting of some areas only, for reactive metallic

surface, by sand blasting, for example, attack may be so rapid that uniform dissolution is observed [6]. Tramp materials, such as iron-containing shot blast or silica-containing sandblast cleaning media can be entrained on a casting surface and increase the corrosion rate. Flux inclusions can also result in localized attack, but this problem has been eliminated by the current industry practice of fluxless melting. Stable corrosion pits initiate at flaws adjacent to a fraction of the intermetallic particles present [27] as a result of the breakdown of passivity. This is followed by the formation of an electrolytic cell of which the intermetallic particle is the cathode of the type AlMnFe, Mg17Al12, or Mg2Cu and the surrounding Mg matrix the anode [28]. Hydrogen evolution is the predominant cathodic reaction and where applicable the Fe/Mn ratio within the AlMnFe intermetallics appears to determine the overall corrosion rate. There is no evidence of initiation at particle-free areas, and the resultant surface is very porous. The a-Mg matrix corrodes preferentially leaving the more noble intermetallics in relief along the grain boundaries [29]. The corrosion of Mg–Al alloys in NaCl solutions is characterized by pit initiation and filiform corrosion, which develops into cellular corrosion. Metallographically polished AE alloys with a high Al content exhibited significantly longer induction times for pit initiation than AS, AM, and AZ alloys (ASTM designations, see Section A3.1. at open-circuit condition in the 5% NaCl solution. Pit initiation and growth normally occurred within 1 h of immersion on alloys such as AS41, AM80, and AZ91, whereas AE41, AE42, and AE46 exhibited induction times of a few hours and AE81 more than 24 h. The AZ-, AS-, and AM-type alloys maintain a bright and shiny appearance in the unattacked part of the corroded surface, whereas the AE alloys tend to become dull due to buildup of a relatively thick hydroxide film and formation of numerous small pits, only a few micrometers in depth. The corresponding backscattered electron image (BEI) and X-ray maps indicate a high chloride concentration in the pits and high aluminum concentration in the unpitted areas. Figure 58.9 shows the corroded surface of alloy AE81 after the hydroxide film has been stripped off in chromic acid. The grain bodies with a low Al concentration (location B) corrode at a faster rate than the Al-rich regions along the grain boundaries (location A), as can also be seen in other Mg–Al alloys. However, on AE alloys, the pits do not easily penetrate the Al-rich zones. Good pitting resistance of the die-cast AE alloys is, therefore, attributed to the presence of these Al-rich zones, which appear to act as barriers against pit propagation. If these barriers are removed by homogenization heat treatment, the corrosion resistance is reduced. Homogenized AE81 exhibited corrosion rates > 100 times higher than the As cast material during a 3-day immersion test in 5% NaCl solution. It is not yet clear whether this unusual sensitivity of corrosion to heat treatment is related to the absence of Mn in


FIGURE 58.9. Morphology of corroded AE81 after removal of the hydroxide film. The grain boundaries with Al-rich areas are more resistant than the Al-lean grain [30]. (Reprinted with the permission from SAE paper No. 930755 1993, Society of Automotive Engineers, Inc.)

this AE81 alloy. The corrosion rate of alloys AM80 and AZ91 were only moderately influenced by a similar heat treatment. In general, homogenized specimens exhibited deeper localized attack than the As cast material [30]. The few studies of pitting of Mg and Mg alloys have been concerned with comparing the pitting behavior of cast to that of rapidly solidified Mg alloys. In these studies, two parameters indicative of pitting resistance were measured: (a) ip, the passive current density, which is a measure of the protective quality of the passive film, and (b) Eb, the breakdown potential, which indicates the resistance to the breakdown of the passive film that results in pitting attack. The more positive the value of Eb, the more protective the film on the metal surface. Makar and Kruger [31] showed that rapidly solidified AZ61 (Mg–6 Al–1 Zn) exhibited a breakdown potential that was 200 mV higher than the value found for cast AZ61 in a buffered carbonate solution (pH 10) containing various levels of Cl (Fig. 58.10); the higher the value of Eb, the greater the resistance to pitting. In a buffered borate solution (pH 9.2) containing various levels of Cl , there was no improvement in the Eb values observed for the rapidly solidified alloy. However, the pits formed at 1 V below Eb were hemispherical, apparently forming at defects in the black film that is observed when the cast AZ61 surface is at 1.5 V saturated calomel electrode (SCE). No small hemispherical pits were found on the rapidly solidified AZ61 [20]. B3.2. Crevice Corrosion. Although a form of attack that occurs at narrow gaps (crevices) appears similar to crevice corrosion, it is somewhat different because the corrosion

823

FIGURE 58.10. Anodic polarization scans for cast and rapidly solidified AZ61 (Mg–6Al–1Zn) in pH 10 sodium carbonate–sodium bicarbonate solution with 100 ppm NaCl [31].

observed is caused by the retention in the crevice of moisture, which, being unable to evaporate, promotes the corrosion of the metal in the narrow recess over extended periods. True crevice corrosion is caused by the development of an anodic region within the crevice because of the exclusion of oxygen and a cathode region outside the crevice where the oxygen concentration is high. Corrosion of magnesium is relatively insensitive to oxygen concentration differences [20]. Corrosion in crevices between Magnox A (Mg–0.18 Al) and mild steel, and between Magnox A and Polytetrafluoroethylene (PTFE) occurred in 200 g/m3 NaOH (pH > 11.5) if the Cl concentration was 1 g/m3 or more [32]. B3.3. Filiform Corrosion. Filiform corrosion is typically associated with metal surfaces having an applied protective coating [33]. Its occurrence on bare Mg–Al alloys indicates that highly resistant oxide films can be naturally formed [34]. Filiform corrosion does not occur on bare pure Mg, indicating the strong influence of alloying elements on corrosion products and behavior. The overall variables of significance are temperature, material structure, and polarization of the microgalvanic cell [29]. A diagram showing the mechanism and the products of the filiform corrosion cell of magnesium is presented in Figure 58.11 [35]. After the initiation period of corrosion pits, filiform corrosion dominates the morphology as narrow semicylindrical corrosion filaments project from the pit [36]. Radial propagation is at a much slower rate than that of the filament tips projecting outward. Lunder et al. [37] observed that propagation of the filaments occurs with voluminous gas evolution at the head while the body immediately behind passivates. Electrochemical transport of chloride ions to the head of the filament appears to be an essential component as is precipitation of insoluble Mg(OH)2 by the anodic reaction

824


FIGURE 58.11. Diagram of the filiform corrosion cell in magnesium. Corrosion products and predominant reactions are identified. Filiform corrosion is a differential aeration cell driven by differences in oxygen concentration between the head and tail sections about 0.1–0.2 V. [35] (Reproduced with permission from ASM International, Materials Park, OH.)

with Mg2 þ ions elsewhere along the filament. The corrosion products may vary because they depend on the environment. Filiform corrosion initiates and then develops into cellular or pitting corrosion. Cellular corrosion occurs when a primary initiation site and secondary pits, formed along the filiform corrosion filaments, coalesce to form a corrosion cell with an epicenter at about the original pit initiation site. The growth rate is at a steady radial rate independent of the material temper. Cellular corrosion continues until the cells impinge on one another, at which point they terminate, thereby forming clearly defined cell boundaries [34]. In the As cast condition, compositional variations orient the growth of filiform corrosion. In homogenized alloys, filiform corrosion propagates transgranularly along crystallographic directions. In Mg–Al alloys, precipitation heat treatment disperses the secondary Mg17–Al12 precipitate, which blocks transgranular propagation of filiform corrosion, thereby reducing the corrosion rate [19]. B3.4. Granular Corrosion. Intergranular corrosion of magnesium alloys does not occur because the grain boundary constituent is invariably cathodic to the grain body. Corrosion of magnesium alloys is concentrated on the grains, and the grain boundary constituent is not only more resistant to attack, but is cathodically protected by the neighboring grain. B3.5. Stress Corrosion Cracking. Pure magnesium is not susceptible to stress corrosion cracking (SCC). The Mg–Al alloys have the greatest SCC susceptibility of all the magnesium alloys, and susceptibility increases with increasing aluminum content. The Mg–Zn alloys have intermediate susceptibility, and the alloys that contain neither aluminum nor zinc are the most SCC resistant. No special heat treatments have been found that will reduce or eliminate SCC [38]. Failures of wrought AZ80 aircraft components resulted from excessive assembly and residual stresses [39–40]. Stress sources likely to promote cracking are weldments and inserts. Welded structures of these alloys require

stress-relief annealing. Magnesium castings have been shown to fail in laboratory tests under tensile loads as low as 50% of yield strength in environments causing negligible general corrosion. The apparent low incidence of SCC service failures of castings is attributable to low stresses actually applied or to stress relaxation by yielding or creep when a fixed deflection is imposed. Although laboratory tests are useful in encouraging conservative design of magnesium alloy structures, results of long-term atmospheric tests of tensile-loaded specimens are considered to be very important. Short-term accelerated tests, such as sodium chloride/potassium chromate (NaCl/ K2CrO4) tests, do not predict SCC behavior reliably in practice [23]. The SCC in magnesium alloys is usually transgranular with significant secondary cracking (branching). Initiation of these cracks has been found to occur invariably at corrosion pits. Mixed transgranular and intergranular crack propagation, and occasionally totally intergranular cracking, have also been observed during magnesium SCC (Fig. 58.12) [41]. Inhibition, by nitrate or carbonate ions, of SCC in salt–chromate solutions is believed to be associated with the formation of a stronger, more stable, or more readily repaired passive film [38]. The SCC can also occur in many other dilute aqueous solutions, including the following, in order of decreasing severity: NaBr, Na2SO4, NaCl, NaNO3, Na2CO3, NaC2H3O2, NaF, and Na2HPO4. SCC has also been reported in dilute solutions of KF, KHF2, HF, KCl, CsCl, Nal, Kl, MgCO3, NaOH, and H2SO4, HNO3, and HCl acids. When pH is > 12, magnesium alloys become very resistant to SCC. Increasing temperature accelerates SCC susceptibility of magnesium alloys, but also improves passivation. Creep deformation could improve SCC resistance. Cathodic polarization has been found to reduce SCC in many studies. Anodic potentials increase SCC susceptibility. Fairman and Bray [42] showed that high anodic potentials, which can


825

FIGURE 58.12. The SCC in an extruded Mg–6Al–lZn alloy tested in a salt–chromate solution, showing intergranular crack (a) in the furnace-cooled alloy and transgranular propagation (b) in the water-quenched material [41]. (Reproduced with permission from ASM Transactions, Pittsburgh, PA)

produce a passive film in a single phase, Mg–Al alloys prevent SCC. The Mg17Al12 phase promoted pitting and SCC in a multiphase alloy. It has been recommended that constant stresses applied for long periods of time should be limited to 30–50% of the yield strength to prevent SCC of magnesium alloys in normal atmospheric environments. It has been suggested that the SCC threshold stress is associated with the onset of plastic deformation (i.e., the elastic limit). The 30% yield strength limit recommended for die cast alloy AZ91 correlates with the elastic limit of this material, reported to be approximately one-third of the yield strength [42]. Dissolution models used to interpret transgranular and intergranular SCC include preferential attack, film rupture, or tunneling in specific dissolution processes. Pickering and Swann [43] have proposed corrosion tunnels in certain SCC systems. A mechanically weak, tubular pitted surface is produced along active slip planes. It has been proposed that the role of corrosion is to produce pits or other stress concentrations that cause cracking by cleavage processes, and to remove obstacles that stop the crack. Fairman and Bray [44] proposed that the passage of dislocations on slip planes rupture the surface film, allowing a corrosion pit to develop, which then initiates cleavage. Liu [45] first suggested that cathodically generated hydrogen be related to magnesium SCC. Experimental evidence supported this model, strengthened by the fact that SCC occurs at crack velocities at which only absorbed hydrogen should be present at the tip [46]. A weak, stress-induced magnesium hydride may form and has been observed on the surface of magnesium SCC fracture [38]. Prevention of SCC is based on avoiding alloys that are susceptible to SCC and environments that cause SCC, and on maintaining the stress below the threshold stress for SCC to occur.

Recommendations to avoid SCC [47] are as follows: 1. The constant stress must be below a threshold level reported to be 30–50% of the tensile yield strength [48]. 2. It has been recommended that inserts with a wall thickness greater than 1.25 mm (0.050 in.) be preheated before casting because cast-in inserts may cause SCC due to local residual stresses created in the surrounding magnesium [48]. 3. Bolted or riveted joints can also produce high local stresses that can cause SCC, so that attention should be given to proper joint design and construction. Examples include the use of preformed parts, avoiding overtorquing of bolts, and providing adequate spacing and edge margins for rivets [48]. 4. Tensile residual stresses from welding were found to be particularly dangerous and, as a result, a low-temperature thermal stress relief treatment has become a recommended practice for welded assemblies. 5. Shot peening and other mechanical processes that create compressive surface residual stresses may also be effective in increasing SCC resistance [49]. 6. Cathodic polarization may reduce, or even prevent, SCC of magnesium alloys in aqueous solutions. 7. Coatings have been shown to extend life, but not to totally prevent SCC, with breaks in the coating reducing protection [49]. In one laboratory study, an inorganic coating was found to accelerate SCC of a SCC-resistant alloy under certain conditions [50]. 8. Cladding of a susceptible magnesium alloy with a SCC-resistant sheet alloy. B4. Corrosion Fatigue There is no endurance limit for magnesium and its alloys in fatigue under corrosive conditions, and the slope of the

826


fatigue curve varies with the corrosive environment and the alloy composition. The Mg–l.5%Mn and Mg–2%Mn–0.5% Ce alloys are more resistant to corrosion fatigue than alloys containing aluminum and zinc; also, media such as 3% NaCl or seawater produce a much more rapid drop in the fatigue curve than does tap water [2]. Substantial reductions in fatigue strength of magnesium alloys are shown in laboratory tests using NaCl spray or drops. Such tests are useful for comparing alloys and heat treatments. Figures 58.13 and 58.14 show data obtained on 1.6 mm (0.064 in.) sheet alloys Mg–6%Al–l%Zn–0.2%Mn and Mg–3%Al–l%Zn–0.3%Mn tested with plate-type bending fatigue equipment in a chloride-containing spray of 0.01% NaCl. Figure 58.13 also shows data obtained on protected and unprotected sheet to determine the effect of normal laboratory exposure. The two rates of spray shown in Figure 58.13 produced the same decrease in fatigue strength.

Both alloys had approximately the same susceptibility to fatigue under corrosive conditions. Unprotected metal in the laboratory atmosphere had slightly lower fatigue strength than when protected. Coatings that exclude the corrosive environment are considerated to provide the primary defense against corrosion fatigue [23]. The corrosion environment was significantly detrimental relative to the air environment. Quasicleavage fatigue crack growth mechanisms have been identified in corrosion fatigue of AZ91E-T6 cast magnesium alloy in both air and 3.5% NaCl. Final fracture regions of samples in both environments were predominantly quasicleavage with some ductile dimples [52]. Under fatigue loading conditions, microcrack initiation in Mg alloys is related to slip in preferentially oriented grains. Quasicleavage usually occurs in the initial stages of fatigue crack growth, which is common for hexagonal close-packed

FIGURE 58.13. Effect of spray intensity of 0.01% sodium chloride on the resistance to fatigue of precipitated Mg–6%Al–l%Zn–0.2%Mn sheet [51]. Specimen size—plate-type specimen 1.6 mm (0.064 in.) thick. Surface preparation–aloxite ground. Temperature–about 30 C (90 F).

CORROSION CHARACTERISTICS OF MAGNESIUM ALLOY SYSTEMS

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FIGURE 58.14. Effect of 0.01% NaCl spray on the resistance to fatigue of precipitated Mg–3%Al–l %Zn–0.3%Mn sheet [51]. Specimen size—plate-type specimen 1.6 mm (0.064 in.) thick. Surface preparation—aloxite ground. Temperature—about 30 C (90 F).

cells. Further crack growth micromechanisms can be brittle or ductile and trans- or intergranular, depending on the metallurgical structure and environmental influence [53]. In general, reduction of temperature increases the fatigue life of Mg alloys mainly by lengthening the crack initiation period [54]. Rotating bending fatigue strengths were reduced by 50% in fretting conditions compared to those in air [51]. Oxides and nitrides are often formed on the surfaces of Mg parts subjected to fretting fatigue condition. Surface rolling, sandblasting, or shot peening can reduce fretting. C. CORROSION CHARACTERISTICS OF MAGNESIUM ALLOY SYSTEMS Each group of alloys has its characteristic corrosion behavior that results from metallurgical properties or presence of

certain intermetallics. The general forms and types of corrosion and the specific properties of every alloy should be considered in developing a corrosion prevention strategy. The following groups of alloys can then be identified: A. Zirconium-free casting alloys 1. Magnesium–aluminum alloys 2. Magnesium–zinc B. Zirconium–containing casting alloys 3. 4. 5. 6. 7.

Mg–Zn–Zr alloys Mg–Re alloys Mg–Th alloys Mg–Ag alloys Wrought alloys, with the same divisions as cast alloys 8. Novel alloys

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High-purity alloys are a necessity in the more severe conditions of corrosion, such as immersion in salt solutions; however, in industrial atmospheres, there is little difference in corrosion performance. Saltwater corrosion studies are typically conducted in 3–5% sodium chloride solutions, following ASTM standards G 31-72 [55] for immersion and B 117-90 [56] for salt spray testing. In these test methods, a corrosive environment is simulated, as might be encountered in a marine or an automotive application (e.g., from salty road splash). The chloride solutions, even in small amounts, usually break down the thin protective magnesium oxide film. The corrosion rate of chemically pure Mg in salt water is in the range of 0.30 mm/year or 12 mpy (mils per year). The corrosion resistance of commercial Mg alloys does not significantly exceed that of pure Mg. Within the Mg–Al alloy system, given that additional alloying elements are used in conjunction with Al and that tramp elements are present, manipulation of alloy chemistry and microstructure can significantly improve the corrosion behavior of these alloys. Aluminum is a common ingot metallurgy (IM) alloying element typically added in the amounts of 2–9 wt % for strength and increased fluidity. The typical IM Mg–Al microstructure shows a-Mg dendrites surrounded by a two-phase eutectic structure along the grain boundaries. Greater Al contents tend to form a continuous eutectic structure and may precipitate Mg17Al12. Manipulation of Al content and heat treatment to control the precipitation of b, Mg17Al12, particles through the eutectic reaction at 28 wt % Al can be used to produce a variety of microstructures like precipitated lamellar b phase in heat-treated AZ91 [19]. The slower solidification rates for gravity versus pressure die castings cause increased average grain sizes and increased corrosion rates, although the tolerance levels are not changed. The smaller grain size of the die casting product results in a finer dispersion of the detrimental material, thereby minimizing its effect as a cathode for localized corrosion. This effect is very much evident in rapidly solidified materials. For die cast Mg–Al alloys in the AM, AS, AZ, and AE series tested by salt spray and by immersion in 5% NaCl solution, the corrosion rate increases when the Al content decreases below 4%. The AE alloys exhibit a high resistance to localized attack because the Al-rich coring along grain boundaries appears to act as an efficient barrier against pit propagation in these alloys. The Ferich phases are particularly detrimental, but Al–Mn phases with a low Al/Mn ratio may also have a high cathodic current output. The phases Mg2Si and Al4MM (MM ¼ misch metal) appear to be harmless from a corrosion point of view [30].

C1. Effects of Alloying Elements Alloying elements not only enhance the mechanical properties of Mg, but also impart a significant impact on the corrosion behavior of Mg–Al alloys. Alloying elements can form secondary particles, which are noble to the Mg matrix, thereby facilitating corrosion, or enrich the corrosion products, thereby possibly inhibiting the corrosion rate. Thus, the Mg–Al alloy corrosion behavior depends on the distribution of the alloying elements [19]. C1.1. Aluminum and the b Phase. Increasing concentrations of 2–8 wt % Al in die cast MG–Al alloys decrease the corrosion rate as shown in Figure 58.15. Low Al additions, of 2–4 wt %, result in a-Mg dendrites surrounded by the twophase, a þ b, eutectic at grain boundaries, whereas higher additions, 6–9 wt % Al, tend to precipitate distinct b particles along grain boundaries, depending on solidification rates. Surrounding the Al-rich b phase are local concentrations of up to 10 wt % Al as a result of microsegregation during solidification [5]. The increasing presence of b particles, which begin to appear above 2 wt % Al, may cause, in part, the improved corrosion resistance of the higher Al-content alloys. The passivating effect of the Al-rich b phase, Mg17Al12, results in a low corrosion rate over a wide pH range. Auger depth profiling shows that, as the Al component dissolves, a Mg-enriched film forms in an alkaline media, and as the Mg component dissolves, an Al-enriched film forms in neutral and slightly acidic media. The synergistic effect of both components leads to the decreased corrosion rate of the b phase.

FIGURE 58.15. Corrosion rate of Mg–alloy die cast rods immersed in 5% NaCl solution as a function of Al content [30]. (Reprinted with the permission from SAE paper No. 930755 1993, Society of Automotive Engineers, Inc.)

CORROSION CHARACTERISTICS OF MAGNESIUM ALLOY SYSTEMS

Additions of Al by rapid solidification processing results in decreased corrosion rates without precipitation of the b phase. Faster solidification of IM alloys disperses fine Mg17Al12 particles, which increases the corrosion resistance, as does controlled precipitation of the b phase [57]. Increasing Al concentrations have a beneficial effect on the corrosion behavior of Mg–Al alloys, but the specific mechanism depends on the distribution of the Al within the magnesium matrix. In addition to the alloying ingredients are added certain other metals that usually present in small amounts. In the alloys containing aluminum, for example, iron usually amounts to 0.02–0.05%. By special techniques and care in melting, this level can be reduced to about one-tenth of this concentration. Such high-purity alloys have much better resistance to salt water than do those of normal purity, but their corrosion behavior in industrial atmospheres is very similar. Furthermore, the practical value of the higher resistance to corrosion is largely offset when components are used in electrical contact with other more cathodic metals. The effect of a steel bolt, for example, even when it has been zinc or cadmium plated, is much greater at the point of contact than that of the local cathodes in the impure alloys. Galvanic corrosion at joints with other metals is not markedly less in the case of the high-purity alloys. Nevertheless, such alloys have their place, and, when they can be used without other metal attachments, provide better intrinsic resistance to corrosion by seawater than the alloys of normal purity [6]. Rare earths (RE) are typically added to Mg–Al alloys as cerium-based misch metal containing lanthanum, neodymium, and praseodymium. A typical composition of MM is 50% Ce, 25% La, 20% Nd, and 3% Pr. These have very low solubilities in Mg (Ce, 0.09; La, 0.14; Nd, 0.10; and Pr, 0.09 at %) [4] and react with Al to form Al4RE intermetallics [58]. These intermetallics, with their high melting temperature, resist coarsening relative to Mg2Si and provide enhanced creep resistance at higher temperatures. Compositions of solidified phases are given in Table 58.7 [37]. Corrosion behavior is optimized through alloy chemistry, by minimizing the cathodic sites, which evolve hydrogen gas, or by enriching the corrosion product film, which can inhibit hydrogen gas evolution and decrease the corrosion rate. Microstructural enhancements, which refine the microstructure and homogenize the distribution of alloying elements also, disperse potentially deleterious elements, thereby enhancing corrosion resistance [19]. The potentials of intermetallic phases, prepared synthetically from the pure components by controlled solidification procedures, are given in Table 58.7 [30]. C1.2. Effects of Zn and Si Additions. Zinc makes the Mg alloy electrochemically more noble, thereby minimizing the corrosion rate [59]. Silicon is intentionally added to only the AS alloys, to combine with Mg, forming Mg2Si, which

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TABLE 58.7. Corrosion Potentials of Synthetically Prepared Intermetallic Phases after 2 h in Deaerated 5% NaCl Solution Saturated with Mg(OH)2 (pH 10.5)a Compound Al3Fe Al3Fe(Mn) Al6(MnFe) Al6Mn(Fe) Al4MM b-Mn Al8Mn5(Fe) Mg17Al12(p) Al8Mn5 Al4Mn(Fe) Al4Mn Al6Mn Mg2Si Mg 99.99%

Corrosion Potential (V/SHE) 0.50 0.71 0.76 0.86 0.91 0.93 0.96 0.96 1.01 1.16 1.21 1.28 1.41 1.42

Reprinted with permission from SAE, paper No. 930755, 1993, Society of Automotive Engineers, Inc.) a

precipitation strengthens the alloy and is relatively innocuous to the corrosion behavior. The compound Mg2Si has a corrosion potential of – 1.65 (V SCE), close to the – 1.66/V SCE value for pure Mg in 5% NaCl solution saturated with Mg(OH)2, pH 10.5 [30]. C1.3. Tramp Element Tolerance Levels. The elements Fe, Ni, and Cu are common tramp elements picked up during melting, handling, and pouring operations. Their influence can be seen in Figure 58.16 for die cast AZ91 corrosion specimens in which the tramp elements were singularly increased.

FIGURE 58.16. Die cast AZ91 salt spray performance versus tramp element content. (Reprinted with permission from [60], SAE paper no. 850417 1989, Society of Automotive Engineers, Inc.)

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TABLE 58.8. Proposed Tramp Element Tolerance Level for Selected Me–Al Die Casting Alloysa

excess of that needed to render the Fe content ineffective could be detrimental to corrosion resistance.

Alloy

Fe/Mn

Fe(max)

Cu(max)

Ni(max)

C1.5. Summary. In summary [37],

AZ91B AM60B AS41B AE42X1

0.032 0.021 0.010 0.020

0.0050 0.0050 0.0035 0.0050

0.030 0.010 0.020 0.050

0.002 0.002 0.002 0.005

1. The corrosion rate of high-purity die-cast Mg alloys in chloride environment decreases rapidly with increasing aluminum content, up to 4 wt %. Further Al additions, up to 9 %, gives only a modest improvement in the corrosion resistance. During immersion testing, AE alloys exhibit a lower corrosion rate than AS, AM, and AZ alloys with similar Al content. 2. Intermetallic compounds containing more than a few percent iron are detrimental because they function as efficient cathodes. However, binary Al–Mn phases with a low Al/Mn ratio may also exhibit a relatively high cathodic current output, causing an increase in the overall corrosion rate. 3. The high corrosion resistance of the AE alloys appears to be related to the presence of passive Al-rich zones along the grain boundaries, acting as barriers against pit propagation. 4. Alloying with silicon does not have an important influence on the corrosion properties because the Mg2Si phase formed is a poor cathode. 5. The Al4MM phase particles precipitated in AE alloys exhibit a passive behavior and do not affect the corrosion process to a significant extent. A high resistance to localized corrosion is observed for the AE alloys with a high Al content.

a

Reproduced with permission from [61]. Copyright ASTM.

The specific ASTM tolerances are given in Table 58.8 (B94-92) [61]. These are typically the same or lower for ingots, as tramp elements are commonly picked up during the melting and pouring operations. C1.4. Effect of Iron and the Fe/Mn Ratio. The Mg–Fe phase diagram shows a very low solid solubility of Fe in magnesium (9.9 ppm). In the absence of Mn, virtually all the Fe precipitates in magnesium as Al3Fe, which has a highly cathodic corrosion potential (Table 58.7). Within an appropriate medium, Al3Fe acts as an effective cathode, catalyzing the reduction reaction, especially hydrogen evolution, which controls the corrosion reaction. Due to the low solubility of Al3Fe in Mg, increasing additions of Al result in smaller tolerance levels for Fe. Typically, up to 1 wt % of Mn is added to improve corrosion resistance by reducing the potential difference between iron-containing particles and the matrix. Its beneficial effect is attributed to either Mn combining with the Fe and precipitating to the bottom of the crucible and/or reacting with the Fe left in suspension during solidification [19]. The relationship between the Fe/Mn ratio in the AlMnFe phase and the corrosion rate is shown in Figure 58.17. Mn in

C2. Influence of Heat Treatment Heat treatment can drastically alter the size, amount, and distribution of the precipitated (b phase, MgI7Al12, which in turn alters the corrosion behavior of IM Mg–Al alloys. A T4 heat treatment (solution heat treatment only for 16 h at 415 C to homogenize the alloy) increases the corrosion rate slightly compared to the As cast material, as shown in Table 58.9. Aune [57] attributed this increase to the resolution of (b particles and release of elemental Fe. Aune [57] completed two T6 treatments with different aging times and temperatures—the T6 being a T4 followed by a T5 treatment. Both corrosion rates were well below those of samples as cast and TABLE 58.9. Heat-Treated AZ91 Corrosion Ratesa Condition

FIGURE 58.17. Relationship between the Fe/Mn ratio in the AlMnFe phase (up to 1 mm in size) and the corrosion rate [27]. (Reprinted with the permission from NACE International.)

As cast T4 T6 aged @ 120 C T6 aged @ 205 C a


Corrosion Rate (mg/cm2-day) 10 11 6 1

RAPID SOLIDIFICATION

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of samples given a T4 heat treatment, as shown in Table 58.9. The lower temperature aging treatment formed a speckled precipitate of (b particles, whereas the higher temperature treatment formed a discontinuous platelike b precipitate with more surface. In the T6 temper, filiform corrosion follows the same crystalline directionality, but area, and thus, a greater effect to inhibit corrosion usually stops at or close to the grain boundary. Corrosion attacks the Al-depleted region between the b phase lamellae along the grain boundary, but passivates after a few minutes of propagation. The improved corrosion behavior results from the presence of b particles [36]. C2.1. Grain Refinement. Grain refinement increases the overall grain boundary area, thereby optimizing the distribution and minimizing the size of any possible detrimental intermetallics, such as Fe3Al. The traditional grain refinement method in sand casting is to add an inoculent, which facilitates heterogeneous nucleation during solidification. Additions of strontium to Mg–Al alloys result in reduced grain size and a lower corrosion rate that is attributed not only to the reduced grain size, but also to changes in the oxide layer structure and composition and in the electrochemical properties of the phases present [62]. D. RAPID SOLIDIFICATION In rapid solidification technologies, including spray or droplet formation, continuous chill casting and in situ melting, typical cooling rates are in the range of 105–107 C/s [63]. Use of continuous chill casting typically produces a thin ribbon of metal, which is then broken into small particles. Then, as with the material formed by spray or droplet formation, the material is often consolidated and extruded. Improper processing can have a significant impact on corrosion behavior. The “chunk” effect [64] is caused by surface oxides on powder particles that lead to poor bonding within the final product [65]. Localized corrosion along these prior boundary oxides leads to particle-size pits and high corrosion rates [19]. Corrosion rates for atomized RS alloy are comparable to those of cast AZ91D, though those for melt-spun RS alloys are significantly higher because of the “chunk” effect (Table 58.10).

FIGURE 58.18. Corrosion rates (1 mpy 25 mm/year) of rapidly solidified magnesium alloys tested in 3% NaCl at 21 C compared with some commercial cast alloys (Extr ¼ extruded) [67]. (Reproduced with permission from [15].)

Nonequilibrium phases in RS alloys can influence the corrosion behavior, for example, Makar and Kruger [31] noted an increase in protection against pit initiation in rapidly solidified (RS) AZ61 compared to IM cast AZ61. Corrosion resistance is improved because the more homogeneous microstructures tend to disperse elements and particles that normally act as cathodic centers, and because the extended solubility of various elements may shift the electrode potentials of light alloys to more noble values (Fig. 58.18) [66]. Using rapid solidification processing, a number of magnesium alloys have been produced in the form of melt-spun ribbon, which is then usually mechanically ground to powder, sealed in cans, and extruded to produce bars. Alloy EA55RS (Mg–5 Al–5 Zn–5 Nd) is now available commercially. Microstructures of the bulk products consist of fine grains 0.3–5 mm in size and dispersions of compounds, such as Mg17Al12, Al2Ca, Mg3Nd, and Mg12Ce [62]. Tensile strengths may exceed 500 MNm 2, which compares with maximum values of 250–300 MNm 2 for conventionally cast magnesium alloys. Some alloys show improved creep resistance at moderately elevated temperatures, but others undergo accelerated creep deformation [15].

TABLE 58.10. Corrosion Rates of Selected Materials (mpy)a Material (wt %) Mg 7.7 Al 2.9 Zn 6.6 Ce 0.35 Mn Mg 10.1 Al 2.7 Zn 1.4 Y 0.44 Mn Mg l0.2 Al 3.2 Zn 5.8 Ce 2.7 Mn Mg 11.1 Al 2.4 Zn 3.2 Y AZ 91D a


Atomized Chill Cast Cast 50 60 350 430 28

D1. Effect of Aluminum Hehmann et al. [68], experimentally measured the solid solubility extensions of 22 RS Mg alloys with extension factors ranging from 1.5 to 1000X. The RS Mg–Al alloys with a maximum terminal solid solubility of 23.4 wt % have decreasing corrosion rates with increasing Al contents from 10 to 40 wt %.

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TABLE 58.11. The Mg–Al Alloy Corrosion Ratesa Material RS–MgAlZnSiMn RS–MgZnAlY RS–MgZnAlNd AZ91HP–T6

Corrosion Rate (mpy) 15 8 11 82

a


D2. Additional Alloying Elements The RS Mg–Al alloys, as do IM Mg–Al alloys, require further alloying elements to improve mechanical properties. The elements Y, Mn, Nd, and Ce have been identified as beneficial to corrosion resistance, whereas Si, Zn, Ca, and Li have been identified as harmful to corrosion resistance [69]. The Mn acts in the same manner as for the IM alloys, by combining with the Al and Fe to form Al (Mn, Fe) intermetallic [70]. This effect causes the low corrosion rate of RS–MgZnAlSiMn alloy (Table 58.11), despite the presence of detrimental Zn and Si. D3. Rare Earth Effects Rare earth alloying elements (Y, Nd, Ce, Pr) result in corrosion rates much lower than those of commercial AZ91HP-T6 alloy (Table 58.11). These elements form stable intermetallic particles in rapidly solidified RS Mg–Al alloys that, similar to the Mg2Si particles, pin the grain boundaries and result in a refined microstructure. Because of the fast cooling rate, various forms of the intermetallics have been reported such as Mg17Y3, Mg3RE (RE ¼ Ce, Nd, Pr), and Al2Nd. The improved corrosion behavior of these alloys, compared to IM Mg–Al alloys, is attributed to the refined RS microstructure, formation of a protective film on the surface of the RS sample as a result of reaction of the saline solution with the rare earths, and the inertness of the second-phase particles [71].

E. MAGNESIUM FINISHING To improve corrosion performance and/or for decorative purposes, finishing processes of magnesium and magnesium alloys are carried out, including surface preparation, chemical treatment, and coating. Finishing can also include oil application, wax coating, anodizing, electroplating, and painting. The designer should use the best combination of methods to meet the functional need of the treated part. The degree of superficial corrosion that can be tolerated without affecting performance and the severity of the service environment are determining factors in selecting an optimum finish [72]. A thin oil or wax film is commonly used for storage or shipping. Sand-cast parts are treated before and during machining operations until the final treatments. A dry storage

atmosphere is important. Chemical and electrochemical methods are used for conversion of magnesium surfaces, so that a more corrosion inhibiting and less alkaline to slightly acidic film replaces the natural alkaline hydroxide–carbonate film on magnesium. The converted surface is generally more compatible with organic coatings. Selected phosphate treatments can be as effective as chromates, even in severe exposures such as marine atmospheric environments. Chemical treatment is strongly recommended for paint formulations, which are based on resins with low resistance to alkaline media. Common chemical treatments alone do very little in aggressive environments and may be unnecessary in mild environments [73, 74]. E1. Cleaning and Surface Preparation Mechanical cleaning of magnesium alloy products is accomplished by grinding and rough polishing, dry or wet abrasive blast cleaning, wire brushing, and wet barrel or bowl abrading (vibratory finishing). The most frequently used methods of mechanical finishing are barrel tumbling, polishing and buffing, vibratory finishing, fiber brushing, and shot blasting. Chemical cleaning methods for magnesium alloys are vapor degreasing, solvent cleaning, emulsion cleaning, alkaline cleaning, and acid pickling. Acid pickling is required for removal of impurities that are tightly bound to the surface or insoluble in solvents and alkalis. The ferric nitrate pickle deposits an invisible chromium oxide passivating film. The acetic–nitrate and phosphoric acid treatments remove even invisible traces of other metals [75]. During the process of surface protection, corrosion of Mg can occur. Treatment in a boiling dichromate solution (or the equivalent), followed by a slushy oil application, is satisfactory. E2. Chemical and Electrochemical Finishing Treatments Chemical and electrochemical finishing treatments can be used alone to provide short-term protection against corrosion and abrasion during shipment and storage, or as pretreatments for subsequent finishing methods. E2.1. Chrome Pickle and Chrome-Free Phosphate Treatments. Chrome pickle and chrome-free phosphate treatments can be used to provide a base for paint or short-term protection. The steps in a chrome pickle include alkaline cleaning followed by a cold rinse, chrome pickle [180 g/L of Na2Cr2O72H2O and 120–180 g/L HNO3 (sp. gr. 1.42)], holding in air for 5 days, cold rinse and hot rinse. A dichromate seal can be introduced between the cold rinse and hot rinse for better protection. Also, a dichromate treatment can replace the chrome pickle. In a modified chrome pickle treatment, an acid pickle and caustic dip or another acid

MAGNESIUM FINISHING

pickle before the modified chrome pickle solution are used [75]. The modified chrome pickle provides a uniform coating by optimizing the etching and passivating action of the chrome pickle bath and by thorough cleaning and washing. A number of commercial phosphate treatments provide performance that is comparable to the best chromate-based surface treatments, particularly for new, high-purity die-cast alloys [72, 73]. E2.2. Anodic Treatments. Anodized coatings have varying degrees of porosity and must be scaled for use in aggressive chloride media. The coatings can be infused with various polymers to produce special properties, including lubricating properties. Galvanic anodizing is a low-voltage direct-current (dc) treatment that produces a thin black conversion coating, used mainly as a paint base (chemical treatment No. 9) [75]. A source of electric power is not required. Proper galvanic action requires the use of racks, made of stainless steels, Monel or phosphor bronze. The constituents of the anodizing aqueous bath are (NH4)2SO4, Na2Cr2O72H2O, NH4OH, and the operating temperature is between 49 and 60 C. More substantial coatings, of 5–30-mm thickness, require anodic polarization by external current. Paint base chemical treatment No. 17 and HAE treatment are currently used. In each of these treatments, a two-phase, two-layer coating is produced. The first layer is about 5 mm thick with a light green or greenish tan color. This layer is covered by a secondphase heavier coating, about 30 mm thick and dark green in color. The second layer is vitreous, relatively brittle, and highly abrasive. As an example, the electrolyte for HAE is composed of KOH, Al(OH)3, K2F2, Na3PO4, and K2MnO4, and the current density is 1.5–2.5 A/dm2. The terminating potential is 65–70 V after 7–10 min and 80–90 Vafter 60 min for the HAE treatment. This treatment consists of six or seven steps: alkaline clean, cold rinse, anodize, cold rinse, dichromate bifluoride dip (Na2Cr2O72H2O þ NH4 HF2), dry air, and possibly heat humidity aging. A heavier Cr-22 treatment is a high-voltage process that is commercially available but not currently used. The final potential can be 320 or 350–380 V for heavier coatings. These coatings provide excellent corrosion resistance in mild media and protected unpainted parts of the structure when properly sealed. The sealing posttreatment consists of an immersion for 2 min in a solution of sodium silicate (10% by vol) at 85–100 C [75]. For environmental considerations, two proprietary chrome-free anodizing treatments have been introduced. In a particular military application, these treatments were found to be superior to the two current anodizing treatments (HAE and No. 17) [76–79]. Magoxid-coat is formed in a slightly alkaline bath and results in MgAlO4 and other beneficial compounds on the surface. The innermost, or barrier, layer is extremely thin, followed by a middle ceramic oxide, providing the majority of

833

corrosion protection since it is almost nonporous. The outermost portion of the coating is a very porous ceramic layer [76]. E2.3. Cathodic Treatments. Zinc and nickel are the only deposits used commercially as undercoatings upon which other commonly plated metals are deposited. Standard practice for plating magnesium involves surface conditioning, zinc immersion plating (zincate solution), and a cyanide copper strike ( 8 mm), followed by a standard plating process. Copper–nickel–chromium plating systems on magnesium satisfy decorative and protective requirements. For interior and mild exterior environments, especially in marine atmosphere, a pore-free deposit is required for satisfactory corrosion resistance. Porosity in the base metal promotes porosity in the deposit [75]. E3. Organic Finishing For optimum corrosion resistance, a chemical conversion coating or anodizing treatment is required prior to applying the organic finishing system. It has been found that anodized components provide greatly improved corrosion durability if the porosity is sealed with a penetrating resin prior to the application of primer and topcoat [77]. For military applications in the United States, finishing is controlled by specifications [75]. Baking paints are harder and more resistant to attack by solvents and are preferable to air-drying ones for applications in solvents. Primers for magnesium should be based on alkali-resistant vehicles, such as polyvinyl butyral, acrylic, polyurethane, vinyl epoxy, and baked phenolic. Titanium dioxide or zinc chromate pigments are used as inhibitors in these vehicles. Finish coats should be compatible with the primer. Vinyl alkyds are resistant to alkalis; acrylics are resistant to chloride environments; alkyd enamels are used for exterior durability; and epoxies have good abrasion resistance. The following finishes have an increasing temperature resistance in the following order. Vinyls (150 C), epoxies, modified epoxies, epoxy-silicones, and silicones. Although chemical treatment can retard the natural alkali that forms at any point of the painted film, primers (13 mm thick), which contain an alkali-resistant vehicle, recommended. Oils, alkyds, and nitrocellulose are best avoided, except for mild exposure [73]. Finishing by one (13 mm thick) or more coats depends on the corrosive medium and the acceptable corrosion rate. A current example of a paint system can consist of an epoxy-based primer or sealer and an acrylic top coat [73, 74]. With epoxy and epoxy-polyester, electrostatic powder deposition is used successfully on magnesium alloys. Cathodic electrodeposition of a resin from an aqueous emulsion can give a more uniform coverage than spray or dip systems of painting. In a recently developed process, a hydrogen-rich or magnesium hydride layer is created on the magnesium surface by cathodic electric charging in aqueous solution. This compound is a good base for painting [80].

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The following standards on finishing should be considered: ASTM D1732: Standard Practices for Preparation of Magnesium Alloy Surfaces for Painting; acid and alkaline cleaners, dip treatments, and anodizing described [81]. ASTM D2651: Standard Practices for Preparation of Magnesium Alloy Surfaces for Adhesive Bonding [82]. ASTM D1654: Standard Evaluation of Painted or Coated Specimens Subjected to Corrosive Environments [83]. Accelerated laboratory testing aims to reproduce, in a much shorter time than in the field, natural corrosion and degradation processes of the paint system and the substrate without changing the corrosion/degradation mechanisms occurring in practice. Accelerated corrosion testing that can have different testing modes in different corrosive aqueous media (immersion, alternating immersion, and emersion and spray) as well as in atmospheres with different relative humidity percentages are recommended for corrosion evaluation of coated magnesium alloys. It is always recommended to test the metal or alloy without coating as reference for corrosion rate and corrosion form in the same conditions. Physical and mechanical testing methods such as film hardness are best measured by making microhardness indents on a cross section of the film, but a minimum film thickness of 25 mm is required. This can be increased 10-fold, for example, in the case of an anodized surface layer (keronite process) impregnated with an organic topcoat such as e-coat, PTFE, or powder coat [84]. One of the most effective tests for judging the ability of a paint base to sustain its protectiveness after damage is to scribe the test panels with a sharp instrument that penetrates the paint and coating layer, leaving bare substrate exposed. Once scribed, the panels are placed in salt spray, and evaluations are taken at regular intervals to determine how far corrosion has migrated from the scribe line to undamaged areas (ASTM D1654). The alternate intermittent immersion in salt water or salt spray is often used to compare the corrosion resistance of magnesium alloys to each other and to other materials. A test period of 28 days has been found to be more severe than a 5year atmospheric exposure. The humidity, or condensation, test is a simple variation of the salt spray test. Test panels are exposed to a climate with very high relative humidity (RH) and no electrolytes, usually at 40–50 C. Coated panels are exposed so that moisture condenses on the test face [85]. It is admitted frequently that cyclic corrosion testing (CCT) gives the relative corrosion rates, structure, and corrosion morphology quiet close to those seen outdoors, if paint systems of the same type of binder are compared [86]. Simple exposures like prohesion may consist of cyclic events, between salt fog and dry conditions. More sophisticated automotive methods

call for multistep cycles that may incorporate immersion, humidity, condensation along with salt fog and dry-off events. The results from an acid rain test (2 h of spraying with acid rain of pH 3.5 plus 1 h drying) did not distinguish the various coated specimens in terms of corrosion behavior. The prohesion test was more corrosive than the acid rain test [84, 86].

ACKNOWLEDGMENTS This chapter is based on the initial and pioneer contribution of W. S. Loose in the first edition of H. H. Uhlig, Corrosion Handbook. Participation and discussions of Dr. I. Nakatsugawa from the Institute of Magnesium Technology, Inc., Sainte-Foy, Quebec, Canada, are deeply acknowledged.

REFERENCES 1. D. S. Tawil, “Magnesium Technology,” in Proceedings of the Conference of The Institute of Metals, Nov. 3–4, 1986, The Institute of Metals, Ed., London, UK, 1987, pp. 66–74. 2. W. S. Loose, Magnesium and Magnesium Alloys, 1st ed., Corrosion Handbook, H. H. Uhlig (Ed.), sponsored by The Electrochemical Society, Wiley, NY, 1976, pp. 218–251. 3. T. R. Massalski, Binary Phase Diagrams, 2nd ed., Vols. 1–4, ASM International, Metals Park, OH, 1990. 4. A. A. Nayeb-Hashemi and J. B. Clark, “Mg-Si Phase Diagram,” in Binary Alloy Phase Diagram, 2nd ed., T. R. Massalski, H. Okamoto, P. R. Subramanian, and L. Kacprzak (Eds.) ASM, Metals Park, OH, 1988, p. 2547. 5. J. D. Hanawalt, C. E. Nelson, and J. A. Peloubet, Trans AIME, 147, 273 (1942). 6. L. L. Shreir, R. A. Jarman, and G. T. Burstein, “Magnesium and Magnesium Alloys,” Corrosion, 1, 4.98–4.115 (1995). 7. ASTM B275-94 “Standard Practice for Codification of Nonferrous Metals and Alloys, Cast and Wrought,” in Annual Book of ASTM Standards, Vol. 02.02, Aluminum and Magnesium Alloys, ASTM, Philadelphia, PA, 1994, pp. 282–287. 8. I. J. Polmear, Light Alloys: Metallurgy of the Light Metals, 2nd ed., Edward Arnold, London, 1989. 9. R. S. Busk, Magnesium Products Design, Marcel Dekker, New York, 1987. 10. ASTM B296 (reapproved 1990), “Standard Practice for Temper Designations of Magnesium Alloys, Cast and Wrought,” in Annual Book of ASTM standards, Vol. 02.02, Aluminum and Magnesium Alloys, ASTM, Philadelphia, PA, 1994, pp. 288–289. 11. P. F. King, J. Electrochem. Soc, 113, 536–539 (1966). 12. W. Unsworth and J. P. King, in Proc. Conf: Int. Conf.: Magnesium Technology, The Institute of Metals, London, 1986, p. 25. 13. J. P. King, G. A. Fowler, and P. Lyon,in Proc. Conf: Light Weight Alloys for Aerospace Applications II, E. W. Lee and N. J. Kim (Eds.), TMS, Warrendale, PA, 1991, p. 423.

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14. W. Durako and L. Joesten, in Proc. 49th Annual World Magnesium Conference, International Magnesium Association, McLean, VA, 1992, p. 87. 15. I. J. Polmear, Mater. Sci. and Technol., 10, 1 (1994). 16. S. Krishnamurthy, E. Robertson, and F. H. Froes, “Rapidly Solidified Magnesium Alloys Containing Rare Earth Addition,” in Advances in Magnesium Alloys and Composites, Minerals, Metals and Materials Society, H. Paris and W. H. Hunt (Eds.), Warrendale, Pennsylvania, 1988, pp. 77–88. 17. G. L. Makar, I. Kruger, and A. Joshi, “The Effect of Alloying Elements on the Corrosion Resistance of Rapidly Solidified Magnesium Alloy,” in Advances in Magnesium Alloys and Composites, H. Paris and W. H. Hunt (Eds.), The Minerals, Metals and Materials Soc., Warrendale, Pennsylvania, 1988, pp. 105–121. 18. M. R. Bothwell, “Magnesium”, in The Corrosion of Light Metals, H. P. Godard, W. B. Jepson, M. R. Bothwell, and R. Kane (Eds.), Wiley, New York, 1967, pp. 259–311. 19. B. E. Carlson and J. W. Jones, “The Metallurgical Aspects of the Corrosion Behavior of Cast Mg-Al Alloys,” Light Metals Processing and Applications, METSOC Conference, Que´bec, 1993, pp. 833–847. 20. G. L. Makar and J. Kruger, Int. Mater. Rev., 38(3), 38 (1993). 21. M. Pourbaix, Atlas of Electrochemical Equilibria in Aqueous Solutions, NACE International and CEBEL-COR (Centre Beige d’E´tude de la Corrosion), Houston, TX, 1974, p. 141. 22. G. G. Perrault, Electroanal. Chem. Interfac. Electrochem., 51, 107 (1974). 23. A. Froats, T. Kr. Aune, D. Hawke, W. Unsworth, and G. Hillis, “Corrosion of Magnesium and Magnesium Alloys,” in Metals Handbook, 9th ed., Vol. 13, Corrosion, ASM International, Materials Park, OH, 1987, pp. 740–754. 24. D. L. Hawke, J. E. Hillis, and W. Unsworth, “Preventive Practice for Controlling the Galvanic Corrosion of Magnesium Alloys,” Technical Committee, International Magnesium Association, McLean, VA, 1988. 25. D. Hawke, Asbjorn Olsen, Proc., Society of Automotive Engineers SAE, Paper No. 930751, Detroit, MI, 1993, pp. 79–84. 26. M. R. Bothwell, J. Electrochem. Soc, 106, 1021 (1959). 27. O. Lunder, T. Kr. Aune, and K. Nisancioglu, Corrosion 43(5), 291 (1987). 28. A. I. Asphahani and W.L. Silence, “Pitting Corrosion,” in Corrosion, Metals Handbook, Vol. 13, 9th ed., ASM, Metals Park, OH, 1987, p. 113. 29. W. E. Mercer and J. E. Hillis,“The Critical Contaminant Limit and Salt Water Corrosion Performance of Magnesium AE42 Alloy,” Technical Paper No. 920073, Society of Automotive Engineers SAE, Detroit, MI, 1992. 30. O. Lunder, K. Nisancioglu, and R. S. Hansen, “Corrosion of Die Cast Magnesium-Aluminum Alloy,” Congress and Exposition, Paper No. 930755 SAE, Detroit, MI, Feb. 26–Mar. 2, 1993, p. 117–126. 31. G. L. Makar and J. Kruger, J. Electrochem. Soc, 13(2), 414. (1990). 32. C. Kirby, Corros. Sci. 27(6), 567 (1987).

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33. C. Hahin, “Filiform Corrosion,” in Corrosion, Metals Handbook, Vol. 13, 9th ed., ASM, Metals Park, OH, 1987, p. 104. 34. K. Nisancioglu, O. Lunder, and T. Kr. Aune, Corrosion Mechanism of AZ91 Magnesium Alloy, 47th Annual World Mag. Conf., Cannes, France, IMA, 1992, p. 43. 35. S. C, Dexter, “Localized Corrosion,” in Metals Handbook, 9th ed. Vol. 13, Corrosion, ASM International, Metals Part, OH, 1987, p. 106. 36. T. Kr. Aune, O. Lunder, and K. Nisancioglu, Microstruct. Sci., 17, 231 (1989) and [19]. 37. O. Lunder, J. E. Lein, S. M. Hesjevik, T. Kr. Aune, and K. Nisancioglu, “Filiform Corrosion of a Magnesium Alloy,” 1lth Annual Corrosion Congress, Florence, Italy, 1990, pp. 5.255–5.262. 38. W. K. Miller, Stress Corrosion Cracking of Magnesium Alloys: Materials Performance and Evaluation, ASM International, Materials Park, OH, 1991. 39. M. Vialatte, “Study of the SCC Behavior of the Alloy Mg–8% Al,” Symposium on the Engineering Practice to avoid Stress Corrosion Cracking, NATO, Neuilly-sur-Seine, France, 1970, pp. 5.1–5.10. 40. M. O. Spiedel, Metall. Trans. A, 6A, 631 (Apr. 1975). 41. D. K. Priest, F. H. Beck, and M. G. Fontana, ASM Trans, 47, 473 (1955). 42. L. Fairman and H. J. Bray, Br. Corros. J., 6, 170, (July 1971). 43. H. W. Pickering and P. R. Swann, Corrosion. 19(11), 373f (1963). 44. L. Fairman and H. J. Bray, Corros. Sci., 11, 533, (1971). 45. H. W. Liu, J. Basic Eng., 92, 633 (Sep. 1970). 46. S. P. Lynch and P. Trevena, Corrosion, 44(2), 113 (Frb. 1988). 47. W. K. Miller and E. F. Ryntz, SAE Trans, 92, 524 (1983). 48. E. Groshart, “Magnesium, Part 1–The Metal,” Met. Finish. 83, (10), 17–20 (Oct. 1995). 49. M. A. Timonova, “Corrosion Cracking of Magnesium Alloys and Methods of Protection Against It,” in Intercrystalline Corrosion and Corrosion of Metals Under Stress, I. A. Levin (Ed.), translated from Russian, Consultant Bureau, New York, 1962, pp. 263–282. 50. V. B. Yakovlev, L. P. Trutneva, N. I. Isaev, and G. Nemetch, Protection Met., 20(3), 300 (1984). 51. Beck, A, The Technology of Magnesium and Its Alloys, F. A. Hughes, London, 1940, p. 236. 52. R. I. Stephens, C. D. Schrader, D. L. Goodenberger, K. B. Lease, V. V. Ogarevic, and S. N. Perov,“Corrosion Fatigue and Stress Corrosion Cracking of A291E-T6 Cast Magnesium Alloy in 3.5% NaCl Solution,” Technical Paper No. 930752, Society of Automotive Engineers SAE, Detroit, MI, 1993. 53. V. V. Ogarevic and R. I. Stephens, Annu. Rev. Mater. Sci., 20, 141 (1990). 54. V. A. Serdyuk and N. M. Grinberg, Int. J. Fatigue, 5(2), 79 (1983). 55. Standard Practice for Laboratory Immersion Corrosion Testing of Metals, 1992 Annual Book of ASTM Standards, ASTM, Philadelphia, PA, 3.02 (G31-72), 1992, p. 102.

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56. Standard Test Method of Salt Spray (Fog) Testing, 1992 Annual Book of ASTM Standards, ASTM, Philadelphia, PA, 3.02 (B117-90), 1992, p. 20. 57. T. Kr. Aune, “Minimizing Base Metal Corrosion on Magnesium Products. The Effect of Element Distribution (Structure) on Corrosion Behavior,” 40th World Magnesium Conference, IMA, Dayton, OH, June 12–15, 1983, p. 52. 58. L. Y. Wei,“Development of Microstructure in Cast Magnesium Alloys,” Ph.D. Thesis, Chalmers University of Technology, Goteborg, Sweden, 1990. 59. J. E. Hillis and S. O. Shook,“Composition and Performance of an Improved Magnesium AS41 Alloy,” Technical Paper No. 890205, Society of Automotive Engineers SAE, Detroit, MI, Feb. 1989. 60. K. N. Reichek, K. J. Clark, and J. E. Hillis,“Controlling the Salt Water Corrosion Performance of Magnesium AZ91 Alloy,” Technical Paper No. 850417, Society of Automotive Engineers SAE, Detroit, MI, 1985. 61. Standard Specification for Magnesium-Alloy Die Castings, 1992 Annual Book of ASTM Standards, ASTM, Philadelphia, PA, 2.02 (B94-92), Non ferrous Metal Products, 2.02, 1992, p. 55. 62. G. Nussbaum. G. Regazzoni, and H. G. Jestl and, in Proc. SAE International Congress and Exposition, Technical Paper Series 900792, Society of Automotive Engineers SAE, Detroit, MI, Warrendale, PA, 1990. 63. F. Hehmann and H. Jones, Rapid Solidification Processing of Magnesium Alloys, Magnesium Technology, 1986, Institute of Metals, London, UK, 1987, p. 83. 64. A. Joshi, R. D. Adamson, and R. E. Lewis, “Processing and Properties of Rapidly Solidified Mg-Al-Zn-Y (or Ce) Alloys,” in Magnesium Alloys and Their Applications, B. L. Moredike and F. Hehmann (Eds.), DGM Informationsgesellschaft, Oberursel, Germany, 1992, pp. 495–502. 65. A. Garboggini and H. B. McShane, “Structural Mg-Al Alloys Produced by Rapid Solidification,” Magnesium alloys and Their Applications, in B. L. Mordike and F. Hehmann (Eds.), DGM Informationsgesellsehaft, Ober ursel, Germany, 1992, pp. 503–510 and [19]. 66. S. K. Das and C. F. Chang, in ‘Rapidly Solidified Crystalline Alloys.’ S. K. Das, B. H. Kear, and C. M. Adam (Eds.), Metallurgical Soc. of AIME, Warrendale, PA, 1985, pp. 137–156. 67. S. K. Das and C. F. Chang,Rapidly Solidified High Strength Corrosion Resistant Magnesium Base Metal Alloys,” U.S. Patent, 4,853,035 (Aug. 1, 1989) and [15]. 68. F. Hehmann, F. Sommer, and B. Predel, Mater. Sci., A125, 249 (1990). 69. A. Joshi and R. E. Lewis, “Role of RSP on Microstructure and Properties of Magnesium Alloys,” in Advances in Magnesium Alloys and Composites, H. Paris and W. H. Hunt (Eds.), Minerals, Metals and Materials Society, Warrendale, Pennsylvania, USA, 1988, pp. 89–103. 70. C. F. Chang, S. K. Das, D. Raybould, and A. Brown, Met. Powder Pre., 41(4), 302 (1986). 71. C. F. Chang, S. K. Das, D. Raybould, R. L. Bye, and E. V. Limoncelli, “Recent Developments in High Strength PM/RS

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59 NICKEL AND NICKEL ALLOYS D. C. AGARWAL*

AND

N. SRIDHAR

DNV Columbus, Inc., Dublin, Ohio

A. Introduction B. Nickel and nickel-base alloys for corrosion applications C1. Alloy systems C1.1. Nickel C2. Ni–Cu alloys C3. Ni–Fe system C4. Ni–Si system C5. Ni–Mo system (B family of alloys) C6. Ni–Cr–Fe alloys C7. Alloy 600/601/602CA C8. Alloy 800/800H/800HT/45TM C9. Ni–Cr–Fe–Mo–Cu alloys C10. Alloy 825 C11. “G” family alloys—G/G-3/G-30 C12. 6% Mo alloys C13. Alloy 20 C14. Alloy 33 C15. High-performance Ni–Cr–Mo alloys C16. Alloy C (1930s–1965) C17. Alloy 625 (1960s–present) C18. Alloy C-276 (1965–present) C19. Alloy C-4 (1970s–present) C20. Alloy C-22 (1982–present) C21. Alloy 59 (1990–present) C22. Alloy 686 (1993–present) C23. Alloy C-2000 (1995–present) C24. Alloy MAT21 (1998–present) D. Corrosion behavior of “C” alloys D1. Resistance to general corrosion D2. Localized corrosion resistance D3. Thermal stability D4. Applications of the “C” family of alloys E. Precipitation hardenable alloys F. High-temperature alloys G. Bibliography References

A. INTRODUCTION Within the chemical process industry, as well as other industries, the 300 series stainless steels have been and will continue to be the most widely used tonnage material after carbon steel. Materials of construction for the modern chemical process and petrochemical industries must not only resist uniform corrosion caused by various corrodents but must also have sufficient localized corrosion and stress corrosion cracking (SCC) resistance. These industries have to cope with the technical and commercial challenges of rigid environmental regulations, the need to increase production efficiency by utilizing higher temperatures and pressures, and more corrosive catalysts and at the same time possess the necessary versatility to handle varied feedstock and upset conditions. Over the past 100 years, improvements in alloy metallurgy, melting technology, and thermomechanical processing, along with a better fundamental understanding of the role of various alloying elements, have led to new nickel alloys, which not only extend the range of usefulness of existing alloys by overcoming their limitations but are also reliable and cost-effective and have opened new areas of applications. Nickel and nickel alloys have useful resistance to a wide variety of corrosive environments typically encountered in various industrial processes. In many instances, the corrosive conditions are too severe to be handled by other commercially available materials, including stainless and superstainless steels. Nickel by itself is a very versatile corrosion-resistant metal, finding many useful applications in industry. More importantly, its metallurgical compatibility over a considerable composition range with a number of other metals as alloying elements has become the basis for many binary, ternary, and other complex nickel-base alloy systems, having very unique


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and specific corrosion-resistant and high-temperatureresistant properties for handling the modern-day corrosive environments of chemical process, petrochemical, marine, pulp and paper, agrichemicals, oil and gas, heat treat, energy conversion, and many other industries. These alloys are more expensive than the standard 300 series stainless steels due to their higher alloy content and more involved thermomechanical processing and hence are used only when stainless steels are not suitable or when product purity and/or safety considerations are of critical importance. Corrosion depends on the chemical composition, the microstructural features of the alloy as developed during thermomechanical processing, the various reactions occurring at the alloy/environment interface, and the chemical nature of the environment itself. In this chapter, the major nickel alloy systems are discussed, including their major characteristics, the effects of alloying elements, and, most importantly, the strengths, weaknesses, and applications of these systems. A few words on fabrication are also included because an improper fabrication may destroy the corrosion resistance of an otherwise fully satisfactory nickel alloy.

B. NICKEL AND NICKEL-BASE ALLOYS FOR CORROSION APPLICATIONS Nickel-base alloys can be roughly classified into the following family of alloys, each developed for certain corrosion characteristics and mechanical properties:

Alloy System Ni Ni–Cu Ni–Mo Ni–Si (some with Cr) Ni–Fe

Ni–Cr–Fe Ni–Cr–Fe–Mo

Type of Environment Designed to Resist Against Caustic Nonoxidizing halides, caustic Nonoxidizing

Some Major Commercial Alloys Ni-200/201 Alloy 400, K-500 Alloys B-2, B-3, B-4, etc. Alloy D, Lewmet, D-205 Invar, Pernifer, etc.

Highly oxidizing, nonhalide Controlled expansion, nonoxidizing Oxidizing, nonhalide Alloy 600, 800, 690, etc. 825, G-3, G-30, alloy Oxidizing, moderately 33, etc. localized corrosion environments

Ni–Cr–Mo

Oxidizing, highly 625, alloy 276, alloy 22, C-2000, localized corrosion environments alloy 59, etc. X-750, 718, 945, etc. Ni–Cr–Fe– Precipitation Mo–(Al,Ti,Nb) hardenable, environments similar to those without (Al,Ti,Nb) High-temperature X,214, Nimonic, Ni–Cr–Mo–Fe– corrosion, wear HR160, 230, Co–C–Al–Ti– MP-35N, etc. resistance Nb–rare earth metals Ni–Ti Shape memory alloys Nitinol

Commercial nickel production started as early as 1804, and since then Ni-base alloy chemistries have evolved steadily in terms of metallurgical sophistication and their ability to function in a broad range of corrosive environments. The development of Ni–Cr alloys first occurred in the early 1900s at about the same time as stainless steel was being developed. In the 1930s, the Ni–Mo and Ni–Cr–Mo alloys were first commercialized, but they contained high carbon and silicon contents and were not useful in the aswelded condition. The development of the argon oxygen decarburization process in the 1960s and the electroslag remelting process in the early 1980s enabled the manufacturing of low-carbon, low-sulfur Ni–Cr–Mo alloys, such as alloy C-276. Better understanding of the electronic structures of alloys and their effects on the formation of intermetallic phases in these alloys resulted in the development of metallurgically stable variants of these alloys starting in the mid-1970s and further developments have continued through the present. The Ni–Fe–Cr–Mo corrosion-resistant alloys developed initially were all solid-solution-strengthened alloys, which meant that their strength could be increased only by cold working. This limited the shapes and sizes of parts produced from these alloys. In parallel with the development of solid-solution-strengthened corrosion-resistant alloys, the development of precipitation-hardenable alloys occurred to satisfy the needs of the emerging aerospace and power generation industries. However, most of these early precipitation-hardenable alloys, such as alloys X-750 and 718, did not have sufficient corrosion resistance. Oil and gas production at increasingly deeper zones placed a premium on high-strength, thick-section Ni-base alloys for components such as the tubing hanger. In the last decade this has spurred the development of precipitation-hardenable alloys with quite high corrosion resistance, such as alloy 945. This chapter provides a concise description of each alloy class with some corrosion data and application areas.

NICKEL AND NICKEL-BASE ALLOYS FOR CORROSION APPLICATIONS

TABLE 59.1. Alloying Elements

839

Alloying Elements and Their Major Effects in Low-Temperature Corrosion Environments Main Effects

Other Effects

Al

Provides matrix for dissolving other alloying elements; improves thermal stability towards intermetallic phases and fabricability Generally added as deoxidizer

Cr

Improves resistance to oxidizing environments

C Co

Increases strength at high temperatures Increases resistance to wear

Cu

Improves resistance to nonoxidizing environments, such as sulfuric and phosphoric acid; as binary alloying element with Ni can improve seawater corrosion resistance. Generally added to Ni-base alloys to reduce cost and enhance scrap utilization; in binary Ni–Fe alloys provides low thermal expansion characteristics; in Ni–Cr–Mo–Fe alloys, higher Fe can reduce tendency to form long-range ordering Increases resistance to nonoxidizing environments, such as HCl, H2SO4; in combination with Cr enhances resistance to oxidizing halide environments Mainly added as deoxidizer during melting; in levels found in typical Ni-base alloys has no effect on corrosion In combination with Cr and Mo can increase localized corrosion resistance; increases strength

Improved corrosion resistance to nonoxidizing environments over Fe; resistance to alkalis; increases chloride SCC resistance In some precipitation-hardening alloys added to form g 0 phase Increases solubility of nitrogen and carbon; increases susceptibility to precipitation of sigma phase; in combination with Mo, W, and N can improve localized corrosion resistance Increases formation of detrimental grain boundary carbides Sometimes an incidental impurity; in high amounts can reduce SCC resistance Decreases localized corrosion resistance in oxidizing halide environments

Ni

Fe

Mo

Mn N

Nb (Cb)

Generally added to prevent formation of chromium carbides (stabilization) or to provide precipitation hardening; No effect on corrosion except indirectly through carbide control

P

Found as tramp element; Detrimental to corrosion resistance

S

Found as tramp element; detrimental to corrosion resistance

Sn

Generally undesirable element in manufacturing of Ni-base alloys, especially high-temperature alloys Improves resistance to highly oxidizing, nonhalide environments such as concentrated sulfuric acid and red fuming nitric acid; in binary Ni–Si alloys are generally cast and require at least 14% Si; Ni–Cr–Si alloys require much less Si Generally added to prevent formation of chromium carbides (stabilization) Generally added to prevent formation of chromium carbides (stabilization) or to provide precipitation hardening; In binary Ni–Ti alloys provides shape memory effect Similar to Mo in effect, but half as effective because it is twice as heavy as Mo; increases resistance to nonoxidizing environments, such as HCl, H2SO4; in combination with Cr and Mo enhances resistance to oxidizing halide environments

Si

Ta Ti

W

Increases tendency to form detrimental intermetallic phases, such as sigma; the higher the Cr, Mo, and W levels, the lower must the Fe addition must be preserve metallurgical stability Provides solid-solution strengthening; increases tendency to form detrimental intermetallic phases, such as m phase. High Mn levels could destabilize metallurgical structure by precipitating intermetallic phases Higher Cr (at least 25%) is necessary to dissolve N in Ni-base alloys; lower Cr alloys can form deleterious chromium nitrides May lead to knife-line attack near welds through resolution of carbides and precipitation of chromium carbides; for precipitation hardening through formation of g 00 , at least 3% Nb is needed; increases tendency to form detrimental Laves phase Phosphorus segregation to grain boundaries can lead to enhanced hydrogen embrittlement Sometimes added to improve machinability, but typically not in Ni–base alloys

Sometimes Si is added to cause precipitation hardening through formation of Ni3Si; Si is detrimental to metallurgical stability; Si increases carbon activity and results in enhanced carbide formation; Si also increases tendency to form sigma phase in presence of Cr and Mo May lead to knife-line attack near welds through resolution of carbides and precipitation of chromium carbides May lead to knife-line attack near welds through resolution of carbides and precipitation of chromium carbides; enhances precipitation-hardening effect of Al and Nb Provides solid-solution strengthening; increases tendency to form detrimental intermetallic phases, such as m phase

840


TABLE 59.2. Alloying Elements and Their Major Effects in High-Temperature Alloys Alloying Elements

Main Feature

Other Features

Cr

Improves oxidation resistance; detrimental to nitriding and fluorination resistance Improves oxidation, nitriding, sulfidation, and carburizing resistance; detrimental to nonoxidizing chlorination resistance Independently and synergistically with Cr improves oxidation resistance; detrimental to nitriding resistance Improves high-temperature strength improves creep strength, detrimental for oxidation resistance at higher temperatures Behaves similarly to molybdenum Increases short-term creep strength; may be beneficial in carburizing; detrimental to nitriding resistance Improves strength, helps nitridation resistance, beneficial to carburization resistance, oxidation resistance adversely effected Improves age hardening Slight positive effect on high-temperature strength and creep; detrimental to oxidation resistance; increases solubility of nitrogen Reduces rate of sulfur diffusion, hence helps with sulfidation resistance; improves solidsolution resistance; improves solid-solution strength Improves carburization, nitriding, and chlorination resistance; detrimental to sulfidation resistance Improves adherence and spalling resistance of oxide layer and hence improves oxidation, sulfidation, carburization resistance

Improves sulfidation resistance; beneficial to carburization and metal dusting resistance Synergistically acts with chromium to improve high-temperature degradation

Si

Al

Mo

W Nb

C

Ti Mn

Co

Ni

Y & RE (rare earths)

C. ALLOYING ELEMENTS AND THEIR EFFECTS IN NICKEL ALLOYS Alloying elements have complex interactive effects on the corrosion resistance and metallurgical stability of Ni-base alloys. The effects of various elements in alloys that are used for wet corrosion resistance and for high-temperature corrosion are presented in Tables 59.2 and 59.3, respectively. Some of the alloying elements are common to both types of alloys but impart different property characteristics in each type. Some elements may be undesirable for wet-corrosion alloys but beneficial for high-temperature corrosion alloys and vice versa. More detailed information on nickel and nickel alloys is available in the References, including the product literature of the nickel alloy producers. This chapter focuses primarily on the aqueous corrosion alloy systems, with a very brief section on high-temperature alloys and fabrication.

Helps improve sulfidation resistance; improves age-hardening effects Helps with reducing chlorination resistance

Improves age hardening

Detrimental to nitriding resistance

Improves high-temperature mechanical properties

C1. Alloy Systems C1.1. Nickel Alloy UNS No. 200 201

N02200 N02201

Nia

Cu

Fe

Mn

C

Si

S

99.0 min 0.25 0.40 0.35 0.15 0.35 0.01 99.0 min 0.25 0.40 0.35 0.02 0.35 0.01

Note: UNS ¼ Unified Numbering System. a Ni þ Co. All values max unless noted otherwise.

The two main alloys, commercially pure alloy 200 and its low-carbon version alloy 201, have good resistance, at low to moderate temperatures, to corrosion by dilute unaerated solution of the common nonoxidizing mineral acids, such as HCl, H2SO4, or H3PO4. The reason for its good behavior is that the standard reduction potential of nickel is more noble than that of iron and less noble than copper. Because of nickel’s high overpotential for hydrogen evolution, there is

ALLOYING ELEMENTS AND THEIR EFFECTS IN NICKEL ALLOYS

TABLE 59.3.

841

Comparison of Some Ni–Cr–Mo Alloys in Various Boiling Corrosive Environments Uniform Corrosion Rate (mpy)a

Media

C-276

C-22

686

C-2000

59

Fe2(SO4)3 50% H2SO4 þ 42 g/L ASTM 28B, 23% H2SO4 þ 1.2% HCL þ 1% CuCl2 þ 1% FeCl3 11.5% H2SO4 þ 1.2% HCL þ 1% CuCl2 þ 1% Fecl3 10% HNO3 65% HNO3 10% H2SO4 50% H2SO4 1.5% HCl 10% HCl 10% H2SO þ 1% HCl 10% H2SO4 þ 1% HClb

240 55 26 19 750 23 240 11 239 87 41

36 7 4 2 52 18 308 14 392 354 92

103 10 8

27 4

24 4 5 2 40 8 176 3 179 70 3

a b

231

5

1.5

67

To convert to millimeters per year (mm/year) multiply by 0.0254. At 90 C.

no easy discharge of hydrogen from any of the common nonoxidizing acids and a supply of oxygen is necessary for rapid corrosion to occur; hence, nickel can corrode rapidly in nonoxidizing environments in the presence of oxidizing species, such as ferric or cupric ions, nitrates, peroxides, or even oxygen. Nickel’s outstanding corrosion resistance to alkalies has led to its successful use as caustic evaporator tubes. At boiling temperatures and concentration of up to 50% NaOH, the corrosion rate is 85

58 >85 >85 >85

a b

PRE ¼ % Cr þ 3.3 % Mo. Above 85 C the solution chemically breaks down.

849

PRECIPITATION-HARDENABLE ALLOYS

TABLE 59.5. Localized Corrosion Resistance in Green Death Solutiona Alloy

PRE

CPT( C)

CCT( C)

Crevice Depth at 105 C

C-22 C-276 59 686 C-2000

65 59 76 74 76

120 110 >120b >120 110

105 105 110 110 100

0.35 mm 0.035 mm 0.025 mm

TABLE 59.6B. Sensitizationa

Thermal Stability per ASTM G28B After Corrosion Rate (mpy)b

Sensitization (h) C-276 C-22

686

>500 339 >500c 313c

c

c

1 3

c

17 85c

C-2000 59 >500c >500c

4d 4d

At 870 C(1600 F). To convert to millimeters per year (mm/year) multiply by 0.0254. c Alloy C-276, C-22, C-2000, and 686—Heavy pitting attack with grains falling out because of deep intergranular attack. d Alloy 59—no pitting attack. a b

a b

11.5% H2SO4 þ 1.2% HCl þ 1% FeCl3 þ 1% CuCl2. Above 120 C, the Green Death solution chemically breaks down.

Alloy C-276, showing the beneficial effects of its high PRE number. D3. Thermal Stability The superior thermal stability of alloy 59 is shown in Tables 59.6A and 59.6B. The data indicate the detrimental effects of tungsten and copper on the thermal stability of various alloys of the C family. During welding of heavywalled vessels and/or hot forming of heavy-walled materials, thermal stability is very important in maintaining superior corrosion resistance in the as-welded condition. Other corrosion resistance data and information on physical metallurgy, fabricability, and weldability of alloy 59 have been adequately covered elsewhere [26–28].

Between 1966 and 1998, 60,000 tons of alloy C-276 and C-4 has been used in a variety of industries, some of which are listed in Table 59.7. Most of the C family alloys are covered under the appropriate ASTM (American Society for Testing and Materials), AWS (American Welding Society), ASME (American Society of Mechanical Engineers), and NACE (NACE International) MR0175 standards and other national and international standards.

E. PRECIPITATION-HARDENABLE ALLOYS UNS Alloy No.

C

D4. Applications of The “C” Family of Alloys

K-500 N05500 0.2

The C family of alloys has found widespread application in chemical and petrochemical industries producing various chlorinated, fluorinated and other organic chemicals, agrichemicals and pharmaceutical industries producing various biocides, pollution control (FGD of coal-fired power plants, waste water treatment, incinerator scrubbers) [29–32], pulp and paper, oil and gas (sour gas production), marine, and many others.

X-750 N07750 718 No7718 0.05

TABLE 59.6A. Sensitizationa

Thermal Stability per ASTM G-28A after Corrosion Rate (mpy)b

Sensitization (h) C-276 C-22 1 3

686

>500 >500 >500c >500c c

c

C-2000 59

>500 >1000c c

116c 178c

40d 51d

At 870 C (1600 F). To convert to millimeters per year (mm/year) multiply by 0.0254. c Alloys C-276, C-22, 686, and C-2000—heavy pitting attack with grains falling out because of deep intergranular attack. d Alloy 59—no pitting attack. a b

925 625

N09925 0.02 N07716 0.03 max

Cr

Cu Fe

Mo Nb Ni

Other

— Bal 3 Al, 0.6Ti 15.5 — 7.0 — 0.9 Bal 2.5 Ti 18.0 — 19.0 3.0 5.0 Bal 0.4 max Ti 21.0 2.0 28 3.0 — Bal 2.1 Ti 20.5 — Bal 8.25 3.5 61 1.3 Ti 30 1

—

Plus

The Ni–Cu alloy, K Monel or alloy 500, was used extensively in drill collars and other high-strength components However, many investigators have reported cracking due to hydrogen generated by galvanic coupling with steel [33–35]. The addition of Al, Ti, Nb, and Si causes precipitation of coherent phases upon heat treatment, which increases the strength of the alloys. This characteristic is especially important in manufacturing thick-section or complex-shaped components that cannot be strengthened purely by cold working. These alloys are used in the nuclear industry for components such as hold-down springs in fuel assemblies, high-strength pins, and bolts in reactor cores. They are also used in oil and gas applications for tubing hangers and valve components. A larger number of failures for X-750 have been reported than for alloy 718, and single-aging treatments are

850


TABLE 59.7. Major Industries Using Alloy C-276/C-4/C-22/59/686 I. Petroleum Petroleum refining Oils/greases Natural gas processing II. Petrochemical Plastic Synthetic organic fibers Organic intermediates Organic chemicals—chlorinated/fluorinated hydrocarbons Synthetic rubber III. CPI—Chemical process industries Fine chemicals Inorganic chemicals Soaps/detergents Paints Fertilizer—agrichemicals—herbicides/pesticides Adhesives Industrial gases IV. Pollution control FGD Wastewater treatment Incineration Hazardous waste Nuclear fuel reprocessing V. Pulp and paper VI. Marine/seawater VII. Pharmaceuticals VIII. Sour gas/oil and gas production IX. Mining/metallurgical

found to yield better SCC resistance than the typical doubleaging treatment used in the aerospace industry [36, 37]. Alloy 625Plus has been shown to have greater SCC resistance in the aged condition to sour gas environments than X-750 and alloy 718 [38].

F. HIGH-TEMPERATURE ALLOYS The need for high-temperature materials is encountered in a wide variety of modern industries, such as aerospace, metallurgical, chemical, and petrochemical, and in many applications, including glass manufacture, heat treatment, waste incinerators, heat recovery, advanced energy conversion systems, and others. Depending on the condition of the chemical makeup and on temperature, a variety of aggressive corrosive environments are produced, which could be either sulfidizing, carburizing, halogenizing, nitriding, reducing, and oxidizing in nature or a combination thereof. All high-temperature alloys have certain limitations and the optimum choice is often a compromise between mechanical property requirements at maximum temperature of operation and requirements for corrosion resistance in the corrosive species present.

Alloys designed to resist high-temperature corrosion have existed since the beginning of the twentieth century. Generally, high-temperature metal degradation occurs at temperatures >540 C, but there are some cases where it can also occur at somewhat lower temperatures. Carbon steel, the “workhorse” material of construction in many industries, is attacked by H2S >260 C, by oxygen or air > 540 C, and by nitrogen > 980 C. The new technologies of thermal destruction of hazardous and municipal waste, fluidized-bed combustion, coal gasification and chemicals from coal processes, and the use of “dirty feedstock,” such as heavy oil and high sulfur coal, coupled with demands for higher efficiency and tougher environmental regulations, have necessitated the use of higher alloy systems of iron-, nickel-, and cobalt-base alloys. Alloy systems must provide reliable and safe performance in a cost-effective manner but must also have sufficient versatility to resist changing corrosive conditions due to feedstock changes. The property requirements in materials of construction for high-temperature applications can be classified under mechanical and high-temperature corrosion resistance as follows: Mechanical Properties

Corrosion Resistance Properties

High-temperature strength Stress rupture strength Creep strength Fatigue strength Thermal stability Thermal shock resistance Toughness Other specific properties

Oxidation Carburization Nitriding Sulfidation Halogenation Molten salt Liquid metal Ash salt deposit Others

Requirements will vary for different industries, such as aerospace, heat treating, power generation, and chemical/ petrochemical processing. In nickel-base alloys, the major elements for imparting specific properties or a combination of properties are chromium, silicon, aluminum, titanium, molybdenum, cobalt, tungsten, and carbon. Others, such as yttrium and rare earths, niobium, tantalum, and zirconium, play very specific roles in improving certain high-temperature corrosion characteristics. These alloying elements can also be classified as follows:

Protective scale formers Solid-solution strengtheners Age-hardening strengtheners Carbide strengtheners Improved scale adhesion (spallation resistance)

Cr as Cr2O3, Al as Al2O3, Si as SiO2 Mo, W, Nb, Ti, Cr, Co Al þ Ti, Al, Ti, Nb, Ta Cr, Mo, W, Ti, Zr, Ta, Nb Rare earths (La, Ce) Y, Hf, Zr, Ta

REFERENCES

References [10, 39–42] provide detailed information on high-temperature alloy systems. REFERENCES 1. Bulletin CEB-2 “Corrosion Resistance of Nickle and NickleContaining Alloys in Caustic Soda and Other Alkalies,” International Nickle Co., Inc., Huntington, WV, 1973. 2. D. C. Agarwal, U. Heubner, M. Koehler,and W. Herda, Mater. Perform., 33(10), 64 (1994). 3. M. K€ohler, R. Kirchheiner, and F. Stenner, Alloy B-10, “A New Nickel-Based Alloy for Strong Chloride-Containing, Highly Acidic and Oxygen-Deficient Environments,” Paper No. 481, CORROSION/98, NACE International, Houston, TX, 1998. 4. D. C. Agarwal and U. Brill, “Material Degradation Problems in High Temperature Environments (Alloys—Alloying Effects— Solutions),” Industrial Heating, Oct. 1994, pp. 55–60. 5. D. C. Agarwal and H. Klein, Applications and Material Performance, Nickel Cobalt 97, Vol. 4, The Metallurgical Society of the Canadian Institute of Mining, Metallurgy and Petroleum, Montreal, 1997, pp. 115–129. 6. H. J. Grabke, E. M. M€uller-Lorenz, J. Kl€ower, and D. C. Agarwal, “Metal Dusting and Carburization Resistance of Nickel Base Alloys,” Paper No. l39, CORROSION/97, NACE International, Houston, TX, 1997. 7. D. C. Agarwal and U. Brill, “NiCr27FeSiRE: A New Alloy for Waste Incineration,” Paper No. 209, CORROSION/93, NACE International, Houston, TX, 1993. 8. M. K€ohler, U. Heubner, K. W. Eichenhofer, and M. Renner, “Alloy 33, A New Corrosion Resistant Austenitic Material for the Refinery Industry and Related Applications,” Paper No. 338, CORROSION/95, NACE International, Houston, TX, 1995. 9. M. K€ohler, U. Heubner, K. W. Eichenhofer, and M. Renner, “Progress with Alloy 33, a New Corrosion Resistant Chromium-Based Austenitic Material,” Paper No. 428, CORROSION/ 96, NACE International, TX, 1996. 10. W. Z. Friend, Corrosion of Nickel and Nickel Base Alloys, Wiley, New York, 1980. 11. E. D. Weister, Corrosion, 13, 659 (1957). 12. W. A. Luce, Chem. Eng., 61(3), 254 (1954). 13. R. K. Swandly, in Nickel Base Alloys in Corrosion Resistance of Metal Alloys, 2nd ed., F. L. LaQue and H. R. Copson (Eds.), Van Nostrand Reinhold, Inc., New York, U.S., 1963, pp. 515–552. 14. Metals Handbook, 9th ed., Vol. 13, Corrosion, Metals Park, OH, 1987, pp. 641–657. 15. I. Class, H. Gr€afen, and E. Scheil, Z. Metallk, 53, 283 (1962). 16. M. A. Streicher, Corrosion, 19, 272 (1963); 32, 79 (1976). 17. R. B. Leonard, Corrosion, 25, 222 (1969). 18. R. W. Kirchner and F. G. Hodge, Werkst. Korros, 24, 1042 (1973). 19. F. G. Hodge and R. W. Kirchner, Corrosion, 32, 332 (1976).

851

20. D. C. Agarwal and W. R. Herda, “Alloying Effects and Innovations in Nickel Base Alloys for Combating Aqueous Corrosion,” VDM Report No. 23, Krupp VDM, Werdohl, Germany, 1995. 21. V. Yanish, “Corrosion Testing in a Hazardous Waste Incinerator and Waste Heat Boiler,” in Proceedings of the Second International Conference on Heat Resistant Materials, Gatlinburg, TN, 11–14 Sept. 1995, ASM, Metals Park, OH, pp. 655–656. 22. P. E. Manning, A. I. Asphahani, and N. Sridhar, “New Developments in Ni–Cr–Mo Alloys,” Paper No. 21, CORROSION/ 83, NACE, Houston, TX, Apr. 1983. 23. P. E. Manning and A. I. Asphaphani, “Advanced Materials Technologies of Interest to the Process Industries,” presented at ACHEMA 85, an International Meeting on Chemical Engineering, Frankfurt am Main, Germany, June 9–15, 1985. 24. N. Sridhar, J. B. C. Wu, and P. E. Manning, J. Metals, 37(11), 51 (1985). 25. M. Morinaga, N. Yukawa, H. Adachi, and H. Ezaki, “New Phacomp and Its Application to Alloy Design,” Fifth International Symposium Superalloys, Seven Springs, 1984, ASM Metals Park, OH, 1984. 26. R. Kirchheiner, M. K€ ohler, and U. Heubner, “A New Highly Corrosion Resistant Material for the Chemical Process Industry, Flue Gas Desulfurization and Related Applications,” Paper No. 90, CORROSION/90, NACE International, Houston, TX, 1990. 27. D. C. Agarwal, U. Heubner, R. Kirchheiner, and M. Koehler, “Cost Effective Solution to CPI Corrosion Problems with a New Ni-Cr-Mo Alloy,” Paper No. 179, CORROSION/91, NACE International, Houston, TX, 1991. 28. U. Heubner, “Nickel Based Alloys,” in Materials Science and Technology, A Comprehensive Treatment, R. W. Cahn, P. Haasen, and E. J. Kramer (Eds.), VCH Verlagsgesellschaft, Germany, 1996, Chapter 7. 29. D. C. Agarwal, “Alloy Selection Methodology and Experiences of the FGD Industry in Solving Complex Corrosion Problems: The Last 25 Years,” Paper No. 447, CORROSION/96, NACE International, Houston, TX, 1996. 30. W. R. Herda and W. Romer, “Recent Experiences with Alloy 59–UNS N06059 (DIN No. 2.4605) in Waste incineration Plant Construction,” Paper No. 557, CORROSION/95, NACE International, Houston, TX, 1995. 31. VDM Case History No. 1, “The Waste Incineration Thermal Power Plant in Essen Kamap,” Krupp VDM, Werdohl, Germany. 32. D. C. Agarwal and Miles Ford, “FGD Metals and Design Technology: Past Problems/Solutions–Present Status and Future Outlook,” Paper No. 485, CORROSION/98, NACE International, Houston, TX, 1998. 33. J. G. Erlings, H.W. deGroot, and J. F. M. van Roy, “Stress Corrosion Cracking and Hydrogen Embrittlement of High Strength Non-magnetic Alloys in Brine,” Mater. Perform., Oct. 1986, pp. 28–34.

852


34. L. H. Wolfe and M. W. Joosten, “Failures of Nickel/Copper Bolts in Subsea Applications,” SPE Production Eng., Aug. 1988, pp. 382–386. 35. K. D. Efird, “Failure of Monel Ni-Cu-Al Alloy K-500 Bolts in Seawater,” Mater. Perform., Apr. 1985, pp. 37–40. 36. N. Sridhar and G. A. Cragnolino, “Stress Corrosion Cracking of Nickel-Base alloys,” in Stress Corrosion Cracking, R. H. Jones (Ed.), ASM International, Materials Park, OH, 1992. 37. J. Kolts, “Alloy 718 for the Oil and Gas Industry, in Superalloy 718—Metallurgy and applications,” E. Loria (Ed.), The Minerals, Metals, and Materials Society, Pittsburgh, PA 1989, pp. 329–344. 38. R. B. Frank and T. A. deBold, “Properties of Age-Hardenable Corrosion-Resistant Nickel-Based Alloy,” Mater. Perform., Sept. 1988, pp. 59–66. 39. U. Heubner, Nickel Alloys and High Alloy Special Stainless Steels, 1st ed., Krupp VDM GmbH, Werdohl, Germany, 1987. 40. U. Heubner, Nickel Alloys and High Alloy Special Stainless Steels, 2nd ed. Krupp VDM GmbH, Werdoh, Germany, 1987. 41. U. Brill, “High Temperature Alloys and Their Use In Furnace Construction,” VDM Report No. 15, Werdohl, Germany, June 1991. 42. G. Lai, High Temperature Corrosion of Engineering Alloys, ASM International, Materials Park, OH, 1990.

BIBLIOGRAPHY R. K. Swandly, “Nickle Base Alloys,” in Corrosion Resistance of Metal and Alloys, 2nd ed., F. L. LaQue and H. R. Copsen (Eds.), Van Nostrand Reinhold, Inc., New York, U.S., 1963, pp. 515–552. Publications of The Huntington Alloy Products Division, The International Nickle Co., Inc. Huntington, W VA. Bulletin CEB-1, The corrosion resistance of nickel-containing alloys in sulfuric acid and related compounds, 1983. Bulletin CEB-3, Resistance of nickel and high-nickel alloys to corrosion by hydrochloric acid, hydrogen chloride, and chlorine. Bulletin CEB-4, Corrosion resistance of nickel-containing alloys in phosphoric acid. Bulletin CEB-5, Hydrofluoric acid, hydrogen fluoride, and fluorine. Bulletin CEB-6, Corrosion resistance of nickel-containing alloys in organic acids and related compounds. Individual alloy product brochures of various alloys. Hightemperature and corrosion alloys product literature. V. D. M. Krupp, High Temperature and Corrosion Alloys Product Literature, Krupp VDM GmbH, Werdohl, Germany. Haynes High Temperature and Corrosion Alloys Product Literature, Haynes International, Kokomo, IN. Publications of Nickel Development Institute on Nickel Alloys, Nickel Development Institute, Toronto, Canada. ASTM Book of Standards, Vol 02.04—Non-Ferrous Metals; Vol. 03.02—Wear and Erosion, Metal Corrosion, ASTM, West Conshohocken, PA.

60 TIN AND TINPLATE T. P. MURPHY* Campion Hall, University of Oxford, Oxford, UK

A. Introduction B. Tin B1. Corrosion behavior B1.1. Atmospheric attack B1.2. Aqueous attack C. Tinplate C1. Corrosion behavior C1.1. External corrosion C1.2. Internal corrosion C2. Tin–iron alloy D. Consequences of corrosion References

A. INTRODUCTION Tin is among the longest-used metals in the service of humankind. Its origins go back millennia, and its combination to form bronze constitutes one of our earliest essays in metallurgy. It may first have been encountered by accident, as a result of the chance reduction of cassiterite in a charcoal fire. Combined with copper, it forms bronze, a metal that helped lift our species out of its primitive condition. The other component of tinplate, mild steel, is probably the metal most widely used in the modern world. While steel is less venerable than tin and, indeed, than tinplate, its development has shaped our present society. The combination of these two metals has proved valuable for centuries and remains today a material of great economic and environmental importance.

*

Deceased.

The properties and uses of these metals, the manner in which they may corrode and the factors influencing corrosion, and the consequences of corrosion are reviewed in this chapter. Given the major uses of the metals concerned, their interaction with other metals and with some nonmetallic materials, other than corrodants, will, of necessity, form part of the story.

B. TIN Tin is a soft, malleable metal with a high luster. It displays allotopic modification, with two forms, a- and b-tin, having densities of 5.8 and 7.3, respectively. The former, known as gray tin, is a brittle, gray, cubic material with little strength. It is the low-temperature modification and, theoretically, may be formed below about 13 C. In practice, temperatures below zero Celsius are required, and the tin must be very pure. Quite small levels of impurity atoms inhibit the transition, and heating to room temperature restores the b form. This, known as white tin, is the familiar, metallic form, ductile and malleable, with a tetragonal structure. Room temperature represents 60% of its absolute melting point at 232 C, so that the mechanical strength of tin is limited. Uses for the metal in its elemental form are thus restricted. It appears, rather, in the form of alloys or coatings. Some use is made of block tin in the brewing industry, but this has largely been superseded by stainless steel. Pewter, an alloy of tin (>90%) with antimony and copper and, nowadays, no lead, is used in decorative work and drinking vessels, and alloys with (decreasingly) lead, bismuth, and other metals find use in solders. Low-melting tin alloys of quite widely differing compositions have served as diecasting materials, and press tools have been made from fusible tin alloys [1]. Solders and tinplate today form the two major uses for tin, with the former having recently


853

854

TIN AND TINPLATE

overtaken the latter. Replacement of more toxic materials such as lead in, for example, wine capsules and gunshot has recently provided new applications for this metal. B1. Corrosion Behavior B1.1. Atmospheric Attack. Tin is not a noble metal but is not among the most active. It is subject to attack both from the liquid and the vapor phases. The attack depends on the nature of the environment. Tin remains reasonably bright for long periods at ambient temperatures in air free of moisture and pollutants. Oxide formation occurs more quickly as the temperature increases and may produce interference effects. Normal domestic atmospheres have only limited corrosive effects on tin. Pollution changes this, with SO2 and H2S capable of producing darkening. The results of such atmospheric attack tend to be formation of oxides and sulfides; the latter are colored so that the attack is obvious. Loss of structural integrity is not a common outcome of atmospheric attack, though the metal is vulnerable to the halogens and vapors of strong mineral acids, and to conditions that lead to formation of films of liquid on the metal. Solderability and electrical resistance can, however, be affected by the formation of surface oxide films, though this is not usually a problem when the tin is sheltered from the weather. Solders are widely used in the electronics industry, with a neareutectic 60% tin, 40% lead composition chosen for its sharp melting point. Few corrosion problems arise in these sheltered applications, though some care is needed in the choice and use of fluxes, which are, of their nature, corrosive. Corrosion rates have been measured for exposed conditions. Britton [2] quotes the work of Hiers and Minarcik on American Society for Testing and Materials (ASTM) tests on tin specimens exposed for periods of up to 20 years. Average penetration rates in micrometers per year (mm/year) were as follows: Rural Industrial Marine

0.05 0.125–0.175 0.175–0.275

B1.2. Aqueous Attack. The potential–pH (Pourbaix) diagram for tin [3] sets out the regions of stability of the metal and various oxides and hydroxides in the presence of water at 298 K (Fig. 60.1). The diagram does not include data for the metal in the presence of complex-forming materials. We note that tin has no zone of stability corresponding to that of water, so that the formation of oxides and hydroxides is to be expected. The diagram indicates that these are reasonably stable between pH values 3 and 10, so that, in the absence of complexants, attack on the metal should be restricted. Outside this range, general attack is quite likely, with Sn2þ or Sn4þ ions being formed at lower pH values and stannites

or stannates under more alkaline conditions. Within it, complexants or species capable of causing local attack may cause corrosion. Thus, in distilled water, attack on tin is limited to the production of an oxide film. This increases in thickness with time and increased temperature and may or may not be protective. The presence of ions like bicarbonate or borate can lead to reinforcement of the film and enhanced protection; other ions can give rise to film weakness and the possibility of localized attack. Some surfactants fall into this category, and chloride is always a potential source of local attack. The formation of oxyhalide films is possible and these can offer some protection. The decisive factor is the solubility of the salts formed. The protective nature of the tin oxide film may be enhanced by the use of oxidizing agents. In practice, few difficulties are encountered in neutral or nearneutral media with tin or high-tin alloys like pewter. The latter has served for many years as a material for drinking vessels. Tin is then a base metal with a reasonably stable oxide film. It forms complexes with a range of ligands, though the Pourbaix diagram does not show this. These complexes may be critical to the performance of the metal. Solders in contact with soft water may be attacked, though, as reported by Britton [2], attack on the lead appears to occur preferentially. In recent years, high-tin solders have replaced lead-containing systems in domestic heating and water distribution systems on environmental grounds, with few incidences of failure, although some electrochemical evidence exists for attack on solder in heating systems. A range of inhibitors is available for this application, as also for use in automobiles. The work of Mercer [4] and his colleagues at Teddington remains definitive in this area, where multimetal systems are common. Solders are now less common in packaging systems, and for food packaging those remaining in use are high tin. Can failure by dissolution of solder is rare, though certain inhibitors for automotive antifreezes have been known to attack high-lead solder side seams in tinplate containers. Leidheiser [5] also gives a good general treatment. Corrosion can sometimes resemble allotropic transformation (tin pest), as shown in Figure 60.2 Here the corrosion is in the specimen on the “outside” of the curve.

C. TINPLATE As we pass to a consideration of tinplate, one further aspect of the corrosion of tin should be borne in mind. The foregoing sections have dealt largely with the resistance of tin to attack and oxidative action (i.e., anodic processes). Corrosion needs a cathodic process to consume electrons produced in the anodic step. Tin has a high hydrogen overvoltage; the reduction of protons then takes place on tin only with

TINPLATE

855

FIGURE 60.1. Potential–pH (Pourbaix) diagram for tin. (Reproduced from [3] by permission of CEBELCOR/NACE.)

difficulty. Corrosion is likely to be slow in the absence of oxygen or other cathodic depolarizers. Contact with a metal having a low hydrogen overvoltage, such as steel, will tend to accelerate the attack. The initial reactions may be summarized as Sn ! Sn2þ þ 2e O2 þ 2H2 O þ 4e ! 4OH Hþ þ e ! 0:5H2 with the first cathodic reaction proceeding readily on tin, while the latter takes place more readily on a material of lower hydrogen overvoltage. Tinplate itself consists of a sheet of mild, low-carbon steel having controlled levels of copper, phosphorus, and sulfur and coated with a layer of tin. The tin is, nowadays, universally coated by electrodeposition from acid, halogen, or less commonly alkaline baths. Bath chemistry has been the object of much study, and many recipes exist to produce good coatings. Sheet thickness is commonly 0.25 mm, though a range of 0.15–0.5 mm may be encountered; the tendency is to thinner stock. The tin coatings are on the order of 0.4 mm, though they are more commonly expressed in terms of

coating mass. Present values range from about 2 g/m2 to 11 g/m2 on each surface. After plating, the tin undergoes momentary fusion by inductive or resistive heating followed by quenching in water. The matte, as-deposited coating is converted to a bright, reflective state, and a layer of a tin–iron (FeSn2) alloy is formed between the tin and the steel substrate. A layer of oxide arises on the tin, and this is usually modified by a “passivation” treatment. This process, using chromic acid or sodium dichromate solutions, sometimes with imposed current, leaves some chromium species in the surface. The object of “passivation” is to control oxidation, suppress sulfide formation in use, and facilitate lacquer coating. A layer of oil, commonly dioctyl sebacate (DOS), is applied to facilitate handling. Thus, from the inside, we have steel, alloy, tin, oxide with chromium, and oil [6]. Tinplate finds its main application in containers for packaging. These cans may be three piece or two piece. The former are produced by forming rectangular blanks round a mandrel and soldering (now increasingly rare), resistance welding, or occasionally cementing the side seam. This produces a cylindrical body to which two ends stamped from tinplate or another sheet material are seamed.

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TIN AND TINPLATE

C1.1. External Corrosion C1.1.1 Atmospheric Attack. The effects of condensation during, for instance, transport and storage of containers are those outlined by Uhlig half a century ago [9]. This may be paraphrased as “steel plus air plus moisture equals rust.” The formation of a film of moisture is important in the atmospheric corrosion of tinplate; it may arise by condensation or from wetting. In the dry condition, slow growth of oxide is the expected outcome, but if humidity control is inadequate, rust may arise. In this context, the tinplate behaves much as steel. Tin, in these conditions, behaves as a cathode for steel, in accordance with its more noble position in the electromotive series. The only protection it offers is as a physical barrier. Thus, rusting is the expected outcome for neutral or near-neutral films of condensation. Changes in this may arise if the nature of the film is changed. The presence of contaminants may alter the processes. Contamination by, for example, cleaning materials may make the film alkaline; spillage of food products may lower the pH. In either case, given the amphoteric nature of tin, detinning may arise.

FIGURE 60.2. Allotropic modification and corrosion of tin; corrosion occurred in the specimen on the “outside” of the curve. (From [2]; reproduced by permission of iTRI.)

Aluminum or ECCS, a mild steel coated with a mixed chromium/chrome oxide film, may constitute the end stock. Soldering changes the state of the tin and welding modifies the steel near the weld. Both welding and cementing produce lap joints that may have exposed edges and the possibility of crevices. Two-piece (body and end) containers are made by presswork using circular blanks. The shallow cups formed from these are redrawn with clearance between punch, die, and stock (draw–redraw, or DRD) or without (drawing and ironing, DWI). An end is seamed in place after filling. The DRD generates no new surface, and the area of the can is substantially that of the original blank. Prelacquered or plain stock may be used. In DWI canmaking, the surface undergoes major deformation with disruption of the tin coating, and it is normal to apply lacquers after forming. The canmaking processes can change the nature of the can material and thus influence corrosion [7]. C1. Corrosion Behavior External and internal attack on containers may be very different. External corrosion will be discussed first.

C1.1.2 Liquid-Phase Attack. The most obvious examples here would be attack by process or cooling water on the outer surfaces of cans. While many cans are processed in steam at temperatures of about 121 C, some are processed in water at high temperature and pressure. In either case, rapid cooling of the cans to prevent growth of spoilage organisms is needed, and this will normally be in water. Whether static retorts, hydrostats, or continuous cookers are used, the tinplate will be exposed to water, at temperatures from 40 to 100 C, and in the presence of varying concentrations of air [8]. The situation may be aggravated by spillage of product or residues of cleaning agents, as described above, and other water treatments may play a role. Biological control of such waters is of vital importance for reasons of public health. Again, treatments may be employed to deal with hard waters. The agents used for these purposes may aggravate attack, and the interaction between hardness treatments, microbial control, and corrosion and its prevention is one of the more interesting problems associated with the use of water. As water is often used in a “mains-to-waste” mode, in these applications, the use of inhibitors may be limited by economic factors. Breakthrough of chloride from watersoftening systems may, as we might expect, give rise to pitting, and complexation as a result of contamination may accompany pH changes. In general, few problems arise. Careful control of process temperatures and times and of cooling times is needed to ensure product safety; good practice in this area will normally result in the avoidance of corrosion problems. However, poor drying of cans may not only pose microbiological problems but also initiate corrosion which may only develop later. Poor practice regarding the external packaging of the

TINPLATE

cans may make things worse. Shrink wrapping wet cans may trap water with the obvious consequences. Uptake of contaminants from cartons may give problems; in particular, chloride may be leached from the board in contact with the cans and may cause local attack. Spillage of product from a damaged or corroded can may cause attack on other cans and may lead to perforation of these, further spillage, and the loss of the whole pallet. Cathodic disbondment of external lacquers can arise when the cans are in contact with aluminum retort baskets and the latter undergo attack from, say, traces of alkaline cleaning agents. As most food cans are processed externally plain, that is, unlacquered, this is not a widespread problem. Some coated food containers are used, and aerosol containers, which are usually lacquered, are pressure tested in a water bath, giving the possibility of corrosion as for food cans. To summarize, external attack is likely in many ways to resemble that on steel. Rusting, pitting, and perforation are risks for this metal, while detinning will be possible under some circumstances. Of the problems listed above, can perforation is the most serious. In addition to product loss, there exists the risk of ingress of foreign material into the cans, of which the most serious would be microbes in food cans. C1.2. Internal Corrosion. Internal corrosion may be very different from external corrosion. Though the material remains the same, the corrodants and conditions are not. For external attack, we are concerned with, essentially, water modified by contaminants in the presence of air, though retorts are purged to reduce air content. Inside the can, the air supply is limited, and the corrodant media may vary enormously. Food and beverages, domestic cleaners, paints, and decorative materials, industrial products, pharmaceuticals, and aerosol products are among the possibilities [10]. These, in turn, may show great variation. Formulated products may vary according to the functional requirements, materials and commercial factors, regulatory constraints, and market needs. Natural products too may vary greatly. Thus, for example, the type and strain of fruit, where it is grown, sunshine, rain, soil condition, use of fertilizers and pesticides, time of harvest, crop treatment, filling, and processing can all affect the corrosive nature of the product. Internally, then, we are faced with a wide range of potentially corrosive media. Let us examine the most important of these product areas, food and beverages. Of great commercial and humanitarian importance, this area also serves to illustrate the important corrosion characteristics of tinplate. We recall the characteristic corrosion properties of tin as formation of oxides stable over a reasonable pH range, a high hydrogen overvoltage, and the ability to form complexes. These determine tin’s behavior inside the can. The third characteristic, that of forming complexes, changes the electrochemical situation greatly from that

857

obtaining on the exterior of the can. The dissolution reaction presented previously may be modified as follows: Sn !

Sn2þ þ 2e #" þ L Sn2þ ðLÞ

where L represents a suitable ligand such as citrate. Tin ions readily form complexes with the so-called “fruit acids,” such as citric, tartaric, and malic. This produces a diminution in the concentration of free tin ions in solution, which encourages further dissolution and, in accordance with the Nernst equation, depresses the potential of the tin. This becomes sufficiently negative for the “normal” polarity of the tin–iron couple to be reversed, with tin “active” (negative) to the steel. It may thus act as a sacrificial anode. This example of cathodic protection, arising from the complexing action of the natural acids, is vital to the functioning of tinplate cans. The relatively large area of tin exposed in a plain can is an effective protective anode for the small area of steel that is exposed at pores in the tin coating. Should this polarity be reversed, the tin would act as a forcing cathode for attack on the steel. The importance of the complexation has been widely recognized, and many workers have treated this question. In fact, Gouda et al. [11] suggested that this is the most important factor in tinplate corrosion. This view probably underestimates the importance of the initial step, tin oxidation, and the factors controlling this step. In order to form complexes of tin ions, it is helpful to have produced some tin ions. Thus, we must consider what determines the ease with which tin forms its ions. As with all corrosion processes, the electrons produced in the oxidation step must be consumed if the reaction is to continue. The cathodic process that consumes the electrons is often rate controlling, and this is the case with tin dissolution. For the dissolution and subsequent complexation of tin, an effective cathodic depolarizer is needed. As tin has a high overvoltage for the hydrogen evolution reaction (HER), protons are unlikely to form useful depolarizers, at least on the tin. Reduction of hydrogen can proceed easily on steel, and in the absence of other depolarizers it is this reduction on the steel exposed at pores in the tin coating that controls detinning. As tin acts protectively, the iron dissolution is suppressed and loss of tin is the corrosion reaction encountered. In the situation described of the “well-packed can,” in which air has been excluded, the detinning will proceed slowly, and the can will have a useful life of years. As tin dissolves, steel may become exposed, allowing reduction of protons and more rapid detinning, so that the process can become self-accelerating. This is indeed what is observed in practice. As the corrosion progresses, hydrogen gas produced at the steel may build up in the can. This will initially reduce the

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TIN AND TINPLATE

vacuum in the can arising from processing and later cause positive pressure to build up, which eventually causes the can to swell. A swollen can is unsaleable because a swell as a result of hydrogen buildup (a “hydrogen swell”) is not distinguishable in the kitchen from swells produced by gases arising from microbial activity. Thus, the time taken for a can to swell is a measure of its useful life, and an old measure of shelf life was the time taken for 50% of a test batch to swell. Nowadays, the amount of tin dissolved would be a more usual measure. A limit of 200 mg/kg (ppm) seems to enjoy general favor with regulatory bodies, a value consistent with the low oral toxicity of tin. This limit can be reasonably easily met in modern well-packed cans of good quality. Pressure has arisen for much lower limits, perhaps as low as 25 mg/kg, which would be difficult to achieve in a plain can. Where other depolarizers are possible, the situation changes. If cans are packed with too much residual oxygen in the headspace, as a result of poor packing practice, the oxygen can undergo ready reduction on the tin surface leading to rapid loss of tin, exposure of iron, and premature can failure. Some foods or drinks may contain natural depolarizers. Fruit colors such as anthocyanins may be readily reduced on tin with consequences similar to oxygen. In addition, the reductive bleaching of these materials leads to color changes in the product that may be unacceptable. In all cases, control is by way of the cathodic process, controlling either by the depolarizers or the reaction surface. Good packing practice with hot filling and steam injection or vacuum closing can reduce the headspace oxygen to an acceptable level; the residual oxygen is quickly consumed and the reaction rate becomes slow. The choice of good tinplate, free of defects and with a high tin coating mass to reduce steel exposure, helps reduce the risk of proton discharge on steel. Can manufacturing and handling practices that reduce the risk of scratches and dents help reduce steel exposure with the same beneficial effects [10]. Reduction of fruit colors is minimized by the use of lacquered tinplate to reduce the tin area available for reduction. This also reduces the relative tin/iron surface area ratio, with some loss in cathodic protection, but a balance has to be struck and, in practice, lacquered cans are very successful. This happy situation with good cathodic protection of steel by tin is not universal. A number of factors may change it. The relative complexing actions of the fruit acids on tin and iron are not identical. Changes in the acids can influence the corrosion. Thus, in pears, the ratio of citric to malic acids may vary. This may depend on the strain of fruit and upon its degree of ripeness at harvest time. As the citric/malic ratio diminishes, the extent to which the tin potential is negative to that of steel is reduced. This may progress to the stage at which the tin is no longer negative to the steel, and it may in fact become positive. In this case, its cathodic protection

becomes ineffective and its value is as a physical barrier coating on the steel. Should suitable cathodic species be available, the tin may act as a cathode for attack on the steel, causing local dissolution of this metal. Even in well-packed cans without fruit colors, such depolarizers may arise as a result of farming practices. The use of nitrate fertilizers is common, and residues of this may be present in the fruit or vegetables. Nitrate offers a multielectron reduction path, starting with nitrate–nitrite and continuing to hydroxylamine and ammonia. This can readily support anodic processes, and nitrate-induced detinning has been a major area of interest in food packaging with studies at Thionville, Parma, Chipping Campden, and in industrial laboratories worldwide. There does seem to be a threshold pH above which the effect does not occur, but when conditions allow, nitrate can produce rapid loss of tin. It may also, of course, support attack on steel when the tin is a cathode in the system. Pesticide residues may also play a role in steel attack. Dithiocarbamates are known to promote attack on the steel. The mechanism is believed to involve the formation of sulfide species on the surface, causing loss of protection from the tin. This can lead to quite severe attack on the base steel. Some products are themselves likely to attack steel. Thus, Cola-type drinks, having phosphate as well as citrate ions and a low pH, tend to be iron dissolvers. The electrochemical testing of these materials reveals that protection by tin is doubtful at best and that a situation with steel anodic to tin is common. As these products are normally packed in lacquered cans, the potential effects of a large tin cathode are minimized, so that perforation is uncommon. The prior treatment of the steel can be important. with steels that have been worked more being more vulnerable. Thus, doublereduced steel, which has had extra cold reduction, or DWI steel may be more readily corrodible than normally processed material. A further form of corrosion that can affect both metals involves the formation of sulfides. Meats, fish, and some vegetables such as peas contain sulfur-bonded protein species. During processing, these may break down to yield hydrogen sulfide. This readily reacts with both tin and steel. The products are colored, tin sulfide being bluish and iron sulfide black. They are among the less soluble compounds of these metals and offer no food contamination problems. However, they are unsightly and the consumer is likely to reject packs showing this defect. The solution is to use internal lacquers pigmented with zinc carbonate or oxide. The sulfide can react with this, but the resulting zinc compound is off-white and causes no concern to the consumer. C2. Tin–Iron Alloy The tin–iron alloy in tinplate is an interesting material. It appears to be a true intermetallic and is quite brittle. Though

CONSEQUENCES OF CORROSION

some differences have arisen over its composition, the accepted stoichiometry is FeSn2. It forms principally during the “flow melting” or refusion process. Examination of non-flow-melted (NFM), as-plated, matte material reveals that some alloy is formed during the plating process. This is present only to the extent of 1% of the level found on flow-melted material. The alloy tends to adopt a cathodic or positive potential with respect to both tin and steel. It is an effective electrode for proton discharge. Nonetheless, performance of tinplate in packing acid fruits is better with a highly continuous alloy layer at the bottom of the pores in the tin coating. This arises as a result of its impermeability to atomic hydrogen. The latter is formed during proton reduction and is known to diffuse readily through mild steel. The higher the rate of diffusion, the faster is the potential corrosion. Tin offers a barrier to hydrogen diffusion; so also does FeSn2, hence the value of a continuous alloy coating. Some of the empirical “special property tests” developed by industry recognize this. Thus the alloy tin couple (ATC) test and the aerated medium polarization (AMP) test sought to assess the resistance of the alloy. Serious deformation, as in DWI canmaking, of course, destroys the brittle alloy.

D. CONSEQUENCES OF CORROSION Two consequences are of major significance: product contamination and package integrity. The first relates, of course, to metal pick-up. Tin and its simple inorganic compounds are, for practical purposes, nontoxic. Cats fed orange juice with some thousands of milligrams per kilogram tin become sick, and human volunteers encounter gastrointestinal upset at some hundreds of milligrams per kilogram. Since 1890, there have been only a few well-documented cases of intoxication by intake of tin in food. All involved thousands of milligrams per kilogram tin, and in each case the problem had arisen by inappropriate use of containers. The effects were, happily, transient in all cases. Nevertheless, a limit of 200 mg/kg tin is applied in most parts of the world. At this level, organoleptic changes become apparent, and visual changes may arise. Tin in beer can give rise to haze, and the color of some foods may be affected. Some foods (e.g., asparagus) benefit from the presence of a little dissolved tin, which preserves the color. Iron is no more toxic than is tin, and no problems of iron poisoning have been reported. Flavor and color changes are, however, a consequence of iron pick-up and most food producers specify low iron content. Thus, iron at a few milligrams per kilogram can affect the taste of soft drinks, and the manufacturers specify subparts per million levels after 6 months storage. In the case of both tin and iron, it is often the commercial demands rather than the regulations that determine the limits, the former frequently being the more rigorous.

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Perforation and loss of package integrity is a more serious matter. The function of the package is to contain the product, protect it from the environment and vice versa, and deliver it to the user, all without costing too much. Perforation of the container, be it through the wall or a seam, allows loss of product. This may, in turn, contaminate or cause attack on other containers. Importantly, perforation may also allow ingress of foreign material, in particular microorganisms. After processing, the contents of a food can are in a state of commercial sterility. This means that they are free of viable forms of microbes having public health significance and of microbes not having public health significance capable of growing and reproducing under normal conditions of ambient storage. Thus, processed foods in intact cans are safe. If microbes gain ingress as a result of perforation, this no longer holds, and a risk of poisoning may arise. This occurrence is very rare. Only one case comes to mind over the last three decades—of pathogenic contamination as a result of corrosion perforation. The internal steel score of meat cans packed with too much air and in a high-chloride environment failed by perforation, allowing ingress of Clostridium botulinum. Normal industrial test procedures detected the situation, and the cans never reached the market. Given the many thousands of millions of food and beverage cans packed each year, the absence of real problems is encouraging. For nonfood products, the risks to life are less apparent; commercial aspects will be important. Loss of product and market may be significant. Contamination both of other products and the environment may arise. The uses of tin and tinplate mean that the risks of catastrophic failure associated with bridges and oil rigs are diminished. However, tinplate has long been used to manufacture automotive brake reservoirs/master cylinders, and failure here could be disastrous. Again, the failure by corrosion of a tin alloy bearing on a marine drive shaft can have serious consequences, especially if the vessel is an oil tanker on a lee shore. These two examples of consequential failure of devices or machines as a result of corrosion of a tinplate or tin component reflect the uses and corrosion characteristics of these materials. For obvious reasons tin finds no structural uses, and tin as a coating on structural steel would offer few benefits. Zinc, cadmium, and such protective coatings are more usual, though tin–zinc alloys have some uses here. The impact of tin corrosion may be out of all proportion to the actual loss of metal. The failure of a small component may give rise to much greater losses. For example, failure of an electronic device in, say, a space probe may vitiate an entire mission. While other causes, for example, tin whiskering, may produce such problems, corrosion and oxidation may play a role. Tin has been described as a “technologist’s metal,” because of its critical use in small quantities. Even the usage in tinplate demands only modest tonnage because of the thin

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coatings used. Nonetheless, the metal has an important impact on our lives. Its corrosion behavior is of great significance, and it is of note that the important function of preserving foods by canning depends to a great extent on the unique corrosion characteristics of tin.

REFERENCES 1. B. T. K. Barry and C. J. Thwaites, Tin, Its Alloys and Compounds, Ellis Horwood, Chichester, 1983. 2. S. C. Britton, Tin Versus Corrosion, iTRI, London, 1975. 3. M. Pourbaix, Atlas d’Equilibres Electrochimiques, GauthierVillars, Paris, 1963, Pergamon, Oxford, 1966. 4. A. D. Mercer, Br. Corros. J., 14(31), 179 (1979).

5. H. Leidheiser, The Corrosion of Tin, Copper and Their Alloys, The Electrochemical Society and Wiley, New York, 1971. 6. W. E. Hoare, E. S. Hedges, and B. T. K. Barry, The Technology of Tinplate, Edward Arnold, London, 1965. 7. E. Morgan, Tinplate and Modern Canmaking Technology, Pergamon, Oxford, 1985. 8. A. Lopez, A Complete Course in Canning, The Canning Trade Inc., Baltimore, MD, 1987. 9. H. H. Uhlig, Corrosion Handbook, The Electrochemical Society and Wiley, New York, 1948. 10. T. P. Murphy, in Progress in the Understanding and Prevention of Corrosion, J. M. Costa and A. D. Mercer (Eds.), Inst. of Materials, London, 1993 pp. 696–711. 11. V. K. Gouda, E. N. Rizkalla, S. Abd-EI-Wahab, and E.M. Ibrahim, Corros. Sci., 21, 1 (1981).

61 TITANIUM AND TITANIUM ALLOYS J. BEEN* AND J. S. GRAUMAN TIMET, Henderson, Nevada

A. B. C. D. E.

Introduction Titanium oxide surfaces General corrosion Pitting corrosion Crevice corrosion E1. Mechanism of crevice corrosion E2. Factors influencing crevice corrosion E3. Detection of crevice corrosion E4. Mitigation of crevice corrosion E5. Specialized forms of crevice corrosion F. Environmentally induced cracking F1. Stress corrosion cracking F2. Corrosion fatigue F3. Hydrogen-induced cracking G. Galvanic corrosion H. Erosion and cavitation resistance I. Microbiologically influenced corrosion/biofouling References

A. INTRODUCTION Titanium metal became a commercial reality in the early 1950s when its high strength/density ratios were especially attractive for aerospace applications. Titanium’s excellent corrosion resistance over a wide range of conditions in many highly corrosive environments has led to a multitude of industrial nonaerospace applications. Some of the first applications in the chemical process industry include wet chlorine gas coolers for chlor-alkali cells, chlorine, and *Present Address: Alberta Innovates Technology Futures, Calgary, Alberta, Canada.

chlorine dioxide bleach equipment in pulp/paper mills and reactor internals for pressure acid leaching of metal ores [1]. In marine environments, titanium is recognized as the best tube material for seawater power plant condensers with almost 650 106 ft installed over the last 40 years and not a single corrosion failure [2]. The market today for titanium is ever expanding. New applications continue to surface as the industry “pushes the process envelope.” The relatively high initial cost of titanium is frequently offset by life-cycle costing, reductions in maintenance and operating cost, and costing on a per-unit area basis as opposed to costing on a per-pound basis. A corrosion allowance is generally not required in designs specifying titanium. Usually the only wall thickness criterion is the pressure or structural requirements for that system. Along these lines, the American Society for Testing and Materials (ASTM) and American Society of Mechaning Engineers (ASME) have recently adopted new higher tensile strengths for unalloyed grades 2, 7, and 16. These new grades, designated 2H, 7H, and 16H, offer a 16% improvement in allowable strengths and thus reduced material weight when designing pressure vessels. All other mechanical and chemical requirements are unchanged from the standard grades 2, 7, and 16, allowing material to be dual certified when being applied to ASME pressure vessels fabrications [2a]. Whereas aerospace applications are mainly concerned with mechanical properties, industrial applications place a greater emphasis on the corrosion resistance of titanium [3–7]. Table 61.1 lists the commercially pure and alloy grades most commonly used in industrial service. Group I contains the commercially pure grades, which differ only in their oxygen and iron content. These grades are highly corrosion resistant, less expensive than titanium alloys, and generally selected when strength is not the main requirement. Increasing oxygen and iron levels improve the


861

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TITANIUM AND TITANIUM ALLOYS

TABLE 61.1. Titanium and Titanium Alloys Commonly Used in Industrial Applications

Common Alloy Designation

UNS Number

ASTM Grade

Nominal Composition (%)

Minimum Tensile Strength (MPa)

Minimum Yield Strength (MPa)

240 345 400 450 550

170 275 275 380 483

345 400 240 345 400 240

275 275 170 275 275 170

620 483 620 620 689

483 345 483 483 586

895 828 828

828 759 759

793 793

759 759

Group I Commercially Pure Titanium Grade 1 Grade 2 Grade 2H Grade 3 Grade 4

R50250 R50400 R50400 R50550 R50700

1 2 2H 3 4

0.06 O 0.12 O 0.12 O 0.2 O 0.3 O

Group II Low Alloy Content Titanium with Pd/Ru Additions Grade 2, Pd Grade 2H, Pd Grade 1, Pd Grade 2, low Pd Grade 2H, low Pd Grade 1, low Pd

R52400 R52400 R52250 R52402 R52402 R52252

7 7H 11 16 16H 17

Ti 3-2.5 Grade 12 Ti 3–2.5, low Pd Ti 3–2.5, Ru Ti 5111

R56320 R53400 R56322 R56323 R55111

9 12 18 28 32

Ti 6–4 Ti 6-4 ELI Ti 6-4 ELI, Ru

R56400 R56407 R56404

5 23 29

0.12 O, 0.15 Pd 0.12 O, 0.15 Pd 0.06 O, 0.15 Pd 0.12 O, 0.05 Pd 0.12 O, 0.15 Pd 0.06 O, 0.05 Pd

Group III Other Alpha and Near-Alpha Alloys 3 Al, 2.5 V 0.3 Mo, 0.8 Ni 3 Al, 2.5 V, 0.05 Pd 3 Al, 2.5 V, 0.1 Ru 5 Al, 1 Sn, 1 Zr, 1 V, 0.8 Mo Group IV Alpha-Beta Alloys 6 Al, 4 V 6 Al, 4 V, 0.13 O max 6 Al, 4 V, 0.1 Ru, 0.13 O max Group V Beta Alloys Beta C Beta 21S

R58640 R58210

19 21

3 Al, 8 V, 6 Cr, 4 Zr, 4 Mo 15 Mo, 3 Al, 2.7 Nb, 0.25 Si

UNS ¼ Unified Numbering System.

material’s strength but reduce its ductility. Whereas Grade 2 can be considered the workhorse of the nonaerospace industry, Grade 1 is selected for applications requiring a high formability. The members of group II offer significantly improved corrosion resistance in reducing media through the presence of small concentrations of palladium or ruthenium. Group III contains other alpha and near-alpha alloys that are characterized by intermediate strength, good ductility, toughness, creep resistance, and weldability. These alloys retain the hexagonal close-packed (hcp) structure characteristic of alpha alloys, which, together with satisfactory strength, make them ideal for cryogenic applications. The presence of molybdenum and nickel in Grade 12, palladium in Grade 18, and ruthenium in Grade 28 improves the corrosion resistance of these grades in reducing acid environments. The alpha–beta alloys of group IV contain an increase in the percent of beta phase and, hence, the strength level as a result of a higher concentration of vanadium, a beta stabilizer. Heat treatments can be used to control the high room temperature strength. Toughness, ductility, and stress corro-

sion cracking (SCC) resistance can be improved by limiting the level of interstitials such as oxygen, nitrogen, and carbon in extra-low interstitial (ELI) content and very-low interstitial (VLI) content grades. Beta alloys are readily cold worked in the solution heat treated and quenched condition. They are heat treatable and can be worked and aged to high strengths at some expense of ductility. Because of their high strength/density ratios, beta alloys are predominantly used in the aerospace industry. An excellent account of physical and mechanical properties of titanium alloys can be found in the Materials Properities Handbook: Titanium Alloys [8].

B. TITANIUM OXIDE SURFACES

Titanium is a reactive metal, ETi=Ti þ 2 ¼ 1:63VSHE [9], owing its excellent corrosion resistance in many environments to a hard, tightly adherent oxide film which forms instantaneously in the presence of an oxygen source.

TITANIUM OXIDE SURFACES

863

FIGURE 61.1. Phase stability diagram of the Ti–H2O system at 25 and 100 C with a titanium ion activity of 10 6 [10].

Figure 61.1 shows the phase stability diagram for the Ti–H2O system [10]. Titanium passivation is present in the stable TiO2H2O area. At higher pH values, HTiO3 dominates in an area that is characterized by corrosion. With increasing temperatures, the area of corrosion extends to lower pH values. The corrosion kinetics are slow within the HTiO3 area as indicated by low measured titanium corrosion rates [7]. At potentials between 1.5 and 2 VSHE, the unstable peroxide TiO32H2O forms [11]. When submerged in a corrosive medium, the overall titanium dissolution rate is very much dependent on the nature and integrity of the oxide. When the oxide is sufficiently thick and stable, electron exchange occurs predominantly with the oxide film. The semiconductive properties of the oxide determine the current/potential behavior of the titanium/oxide system. When the oxide is sufficiently thin (0.4–3 nm), electron exchange occurs between the redox electrolyte and the underlying metal by direct tunneling or resonance tunneling via intermediate states [12, 13]. As a. result of direct tunneling, which consists of electron transfer in one step without loss of energy, electron exchange is under kinetic control with current/potential characteristics that are similar to those of the bulk metal. The oxide functions as a potential energy barrier and the current decreases with increasing oxide thickness. The anodic transfer coefficient becomes smaller with increasing oxide thickness and the cathodic transfer coefficient becomes greater. A cathodic Tafel coefficient of –0.12 V and an anodic Tafel coefficient of 0.12 V yielded electrochemical corrosion rates that compared satisfactorily with weight loss corrosion rates in acid media [14]. The same cathodic Tafel slope but an anodic Tafel slope of 0.25 V [15] yielded good results in alkaline peroxide bleaching environments [16].

A freshly abraded titanium surface immediately passivates to form a crystalline rutile and/or anatase oxide layer. Rutile is the more common titanium dioxide (TiO2) and slightly more stable than anatase by 12 kJ/mol [17–20]. Anatase is a material with the highest photocatalytic detoxification efficiency in ground and surface water purification [21]. Rutile finds application as a catalyst in organic oxidation reactions [22]. The isoelectric point of TiO2 is 6.2, which, together with a high dielectric constant [23], renders titanium oxide waterlike with small electrostatic forces and, consequently, highly compatible as a biomaterial [24]. The titanium oxide gradually decreases in oxygen content from TiO2 at the surface to Ti2O3 and TiO as it approaches the metal oxide interface [25]. Depending on the environment, this oxide may be covered with an amorphous or hydrated surface oxide, giving a two-layer oxide structure. The oxide may be thickened in the presence of oxidizing agents through anodization or thermal oxidation. In reducing acid environments, severe corrosion can be avoided through the application of anodic protection that aids in the formation of a protective surface. For example, the corrosion rate of a titanium heat exchanger in a 40% sulfuric acid environment can be reduced 11,000 times to a rate of 0.005 mm/year (0.2 mpy) through the application of 2.1 Voverpotential [26, 27]. Anodic protection also appears to increase the protective nature of the oxide. Tomashov et al. [28] suggest that this is the result of a decreasing number of defects and decreasing ionic conductivity. Care must be taken not to exceed the repassivation potential of titanium. Anodization at increasingly higher potentials can thicken ˚ to several thousand the very thin natural oxide from 20 A angstroms, depending on the applied potential. As the thickness increases, the oxide progresses through a spectrum of

864


interference colors [29, 30]. Thick oxides were traditionally thought to increase the corrosion resistance. Anodization used to be recommended for heat exchanger tubing to improve crevice corrosion resistance and limit hydriding. Studies later showed that, although there was a slightly higher initial corrosion resistance, the anodized surfaces didn’t behave much better than freshly pickled surfaces in hydrochloric acid solutions [31, 32]. The high dissolution rate of the anodized film has been attributed to the fact that the oxide is amorphous and hydrated [32–34]. Thermal oxidation produces an unhydrated rutile oxide which offers greater corrosion protection than anodized or pickled oxide surfaces [31, 32]. Since the oxide surface is at a more anodic potential, small cracks in the oxide are protected by the anodic corrosion potential in mild reducing acids. As an additional benefit, the thermally oxidized titanium offers a more effective barrier against hydrogen permeation that increases with increasing oxide thickness. The more noble oxide potential may, however, increase the driving force for galvanic corrosion. At elevated temperatures, titanium oxidizes in air to form an oxide scale and an oxygen-rich metal layer. The extent and rate of oxide formation are dependent on the exposure temperature and time. At temperatures below 500 C, the oxidation rate of titanium is low and tends to decrease with time [25, 35]. Long-term exposures at temperatures > 650 C will lead to cracking of the brittle oxide scale and rapid continuous oxide growth [36]. Figure 61.2 illustrates how both the oxide scale and the oxygen-rich metal layer on Grade 2 titanium roughly quadruple in thickness as the temperature is increased from 538 to 649 C. The heat resistance of titanium can be increased through alloying [37] or application of oxidation-resistant coatings [38].

C. GENERAL CORROSION General corrosion is rarely seen in service since titanium is usually not cost effective if a corrosion allowance is necessary. When observed, reducing acids are most often the cause. Titanium offers moderate resistance to mineral acids such as hydrochloric, sulfuric, and phosphoric acid and organic acids like oxalic and sulfamic acid [3, 7]. The corrosion rate varies with acid type, concentration, and temperature. Hydrofluoric acid solutions are routinely used in pickling and etching processes because of the extremely high corrosion rates experienced even at parts per million (ppm) concentrations. Alloying additions of noble metals such as palladium and ruthenium, as well as additions of molybdenum and nickel, were found to be quite effective in increasing the corrosion resistance of titanium in reducing acid environments [25, 39– 41]. Palladium and ruthenium are added in small concentrations, typically 0.05–0.20 wt %, that do not affect the physical and mechanical properties of the titanium alloy but passivate the metal by shifting the corrosion potential into the passive anodic regime [40, 41]. Both nickel and molybdenum reduce the susceptibility of titanium to anodic dissolution. The latter alloying additions increase the alloy strength at the expense of ductility [25, 39]. Small concentrations of oxidizing species effectively increase titanium’s corrosion resistance in reducing acids by positively polarizing the metal. Only parts-per-million concentrations of certain multivalent transition metal ions, nitrates, oxychloro anions, noble metal ions, organic compounds, chlorine, and oxygen are required to induce passivity [3]. Inhibitor levels may be present as contaminants in

FIGURE 61.2. High-temperature oxidation of Grade 2 exposed for 500 h at (a) 538 C and (b) 649 C. The surface oxide layer covers a layer of oxygen-rich alpha structure, accentuated by a lactic acid etch.

GENERAL CORROSION

865

FIGURE 61.3. Effect of ferric ions on the corrosion of (a) Grade 2 and (b) Grade 7 in hydrochloric acid solutions [4]. The isocorrosion plane corresponds to a corrosion rate of 0.13 mm/year.

process streams allowing the safe use of titanium. Figure 61.3 illustrates how hydrochloric acid cleaning solutions can be inhibited by small concentrations of ferric ions. Titanium is known as one of the most corrosion-resistant metals in oxidizing environments as these conditions generally assure oxide film stability. Highly resistant to oxidizing acids such as nitric and chromic acid at room temperature, general corrosion may occur in nitric acid at boiling temperatures in the 20–70 wt % range [42]. When the solution is

not refreshed, buildup of Ti4 þ ions leads to a more protective, dehydrated titanium oxide [43]. Other metal ions, such as Fe3 þ , Ru3 þ , Rh3 þ , Ce4 þ , and Cr6 þ , and oxidizing ions, such as VO2 þ and Cr2O72 , also inhibit high-temperature corrosion [44]. Only tantalum alloying additions significantly improved the corrosion resistance and yielded titanium alloys that were virtually insensitive to changes in temperature and acid concentration [45]. A pyrophoric reaction may develop in red fuming nitric acid following rapid

866


intergranular attack [46, 47]. The presence of sufficient concentrations of water has successfully permitted the long-term use of titanium as a construction material in nitric acid production plants [48]. Titanium is highly corrosion resistant to solutions of chlorites, hypochlorites, chlorates, perchlorates, and chlorine dioxide [3, 7, 49]. Widely used to handle moist chlorine gas, titanium has earned a reputation for outstanding performance in chlor-alkali cells and pulp and paper bleaching equipment. Rapid ignition, forming TiCl4, will occur in dry chlorine. However, a moisture content of 0.4% at room temperature and 1.2% at 175 C is sufficient for passivation [50, 51]. Oxidation of a fresh titanium surface is an exothermic process which may lead to melting of the metal if the heat cannot be removed fast enough. In 12. Dissolved oxidizing species, such as chlorate or nitrate compounds, can be used in alkaline cleaning solutions to extend the resistance to hydrogen uptake to somewhat higher temperatures [3]. The effect of hydrogen peroxide on titanium corrosion has been studied with relation to radioactive waste containers (10 4 M) [63], medical implants (0.01–0.1 M) [64, 65], alkaline pulp, and paper bleaching environments (0.l– 0.2 M) [10, 16] and surface etching or bonding pretreatment [66]. Small concentrations of hydrogen peroxide (1 10 4 M) in brine solutions strengthen the titanium corrosion resistance by shifting the corrosion potential in the noble direction. This shift is attributed to the additional cathodic reduction reaction of hydrogen peroxide and the formation of a thicker oxide layer containing more stable anatase [63]. Addition of a greater hydrogen peroxide concentration, 0.01–0.l M H2O2, still leads to a more passive corrosion potential, but the two-layer oxide becomes increasingly more hydrated and hydroxylated [64, 65]. Further increases in temperature, pH, and peroxide concentration result in increasing corrosion rates and a more active corrosion potential as the oxide thins becomes more conductive and less protective. Small additions of calcium, silicates, and magnesium appear to be effective inhibitors by adsorbing onto the oxide surface and forming a physical barrier to ion transport. Over extended time periods, however, the effect of calcium additions is questionable as corrosion rates begin to increase again [16]. In pulp bleaching solutions, pulp itself proved to be an effective inhibitor under normal operating conditions [16]. Extremely high titanium dissolution rates may be obtained at extreme conditions. A rate as high as 2300 mm/year was measured at 95 C, 2.5 M NaOH, and 0.5 M H2O2 [10]. General corrosion leads to extensive roughening of the surface, which can be attractive for subsequent coating or bonding processes [66]. Titanium is widely used in organic process streams and has been the material of choice for critical areas of terepthalic and adipic acid production. Titanium is highly resistant to solutions of alcohols, aldehydes, esters, ketones, and hydrocarbons [3, 7]. To maintain the integrity of the protective

CREVICE CORROSION

oxide film, some degree of moisture or oxygen should be present. Generally, a moisture content of merely parts per million is sufficient for passivation, a concentration, which experience has shown, can usually be expected in industrial organic processes. Overall, titanium exhibits excellent corrosion resistance in organic acid solutions [3, 7, 67, 68]. Aeration may be required to maintain passivity. Fully resistant in aerated aqueous formic acid, titanium corrosion rates may become unacceptably high at elevated temperatures in deoxygenated formic acid [7]. A strong oxygen effect has also been observed in urea reactors where corrosion potentials dropped as the supply of oxygen stopped. Titanium is a preferred material of construction since it is not affected much by a temporary lack of oxygen [69]. The corrosivity of urea has been linked to an amino–formic acid intermediate [70]. Negative effects of oxygen have been observed in mixtures of anhydrous acetic acid and acetic anhydride, apparently by facilitating the cathodic reduction reaction that favored the production of the acetate ion. Alloying with palladium did not help in this case [71]. Titanium exhibits a poor resistance to corrosion in propionic acid vapor and has a limited stability in oxalic acid solutions. The corrosion rate in the latter medium increases with increasing temperature and acid concentration [67]. The corrosion resistance in oxalic acid can be improved by molybdenum alloying additions [72] or the addition of oxidizing agents such as antimony(III) [73].

D. PITTING CORROSION Titanium exhibits remarkable resistance to pitting attack in chloride media with pitting potentials in excess of þ 5 VSCE in saturated NaCl at boiling. Thus, titanium generally does not exhibit spontaneous pitting under normal circumstances. Pitting resistance is lower in other halide media; however, potentials still remain at or above þ 1 VSCE [74]. Pitting potentials of titanium in sulfate and phosphate media are reported to be in excess of þ 80 VSCE [3]. Alloying can lower pitting potentials somewhat, yet even highly alloyed titanium exhibits pitting potentials greater than þ 1 VSCE in hightemperature NaCl and HCl environments [75]. Pitting corrosion failures of titanium in service are thus extremely rare.

E. CREVICE CORROSION Titanium, being a reactive metal and relying on its passive film for corrosion resistance, is susceptible to localized corrosion in much the same manner as other passive film metals like aluminum, stainless steel, and nickel alloys. Crevice attack can occur on titanium in hot halide or sulfate-containing media. Corrosion can be observed in tight

867

gasket-to-metal or metal-to-metal joints or under adherent deposits formed by a process stream. Under normal circumstances, crevice corrosion resistance will probably be the limiting factor for successful use of titanium; thus a thorough understanding of the mechanism, influencing factors, and mitigation techniques will prove invaluable. Several excellent reference sources are available that compliment the information contained herein [1, 3, 8, 76–78]. E1. Mechanism of Crevice Corrosion As with stainless steels, the mechanism for attack on titanium involves formation of an occluded differential aeration cell, in which slow but finite corrosion depletes the crevice of oxygen through surface oxide formation, as shown in reaction (61.1). Anion migration into the crevice then occurs to preserve mass and charge balance. In the case of chlorides, this results in formation of unstable titanium chloride and oxychloride intermediate compounds that hydrolyze to form free acid, thus lowering the pH in the crevice. At this point, the corrosion reaction becomes selfsustaining as the acid generated from hydrolysis now further attacks the underlying metal. Within the crevice, pH levels of 5 wt % in the a phase, can lead to the formation of Ti3Al, which lowers KISCC and increases the velocity of cracking [94]. The presence of tin further decreases the SCC resistance. Whereas SCC in sea water is not a concern for the lower strength commercially pure grades, susceptibility can be lowered substantially in higher strength aluminum-containing alloys by lowering the oxygen content, as, for example, in the case of Ti 6-4 ELI (Grade 23). Furthermore, in Ti 6-4, an acicular structure provides a lower susceptibility to SCC than an equiaxed morphology, which may be related to the mean free path of the susceptible a phase [94, 95] (see Fig. 61.7). Stress corrosion cracking susceptibility is influenced by the concentration of damaging species, pH, potential, temperature, and viscosity. Addition of halide ions such as Cl , Br , and I– may accelerate or induce SCC, an effect that increases with increasing halide concentration [99]. Reducing the pH results in a greater susceptibility to SCC. Crack velocity was found to increase with increasing temperature and decreasing viscosity. Cathodic protection is effective in neutral aqueous halide solutions at potentials more negative than 1 VSCE but ineffective in acid solutions [99]. Stress corrosion cracking of a and a–b alloys takes place by transgranular cleavage of the a phase, where the a phase controls the overall crack propagation rate. Small differences

871

between the main crack propagation plane and the cleavage planes lead to numerous stepwise facets, the result of lowenergy ductile rupture. These flutes connect cleavage planes and may exceed the grain size [100]. Intergranular corrosion occurs in methanolic halide solutions through the formation of titanium methoxide. As little as 1.5% water is sufficient to hydrolyze the titanium methoxide and passivate the titanium [101]. Higher strength alloys may require the presence of more water depending on product form and alloying content [102]. Noble metal ion additions such as Pd2 þ and Au3 þ facilitate the cathodic process and increase the intergranular corrosion rate. The dissolution process is temperature and viscosity dependent and is accelerated by stress and anodic currents, suggesting a stress-accelerated anodic dissolution or diffusion-controlled mechanism [95]. However, transgranular cracking in medium- to high-strength a and a–b alloys has been attributed to the absorption of hydrogen at the crack tip. Whereas dissolved platinum group metal ions accelerate intergranular corrosion, they have an inhibiting effect on the crack propagation rate, possibly by favoring the recombination of hydrogen atoms as opposed to hydrogen absorption. Since the embrittlement is also dependent on strain rate, it has been suggested that the contribution of stress is, in part, the result of hydrogen embrittlement [103]. A mixed mode of intergranular cracking and transgranular cleavage has been observed in these environments [104]. It has been shown that surface oxide rupture precedes crack initiation with subsequent dissolution and hydrolysis reactions leading to crack tip acidification. The crack tip local pH may be significantly lower than the bulk solution pH.

FIGURE 61.7. A hot salt crack changed to a stress unfavorable direction in Ti 6Al-4V with a Widmanst€atten microstructure [98].

872


A sharp notch or fatigue precrack appears to be required for hydrogen production to occur at the crack tip. Absorption may then occur at a deforming crack tip surface [105]. Hot salt SCC is of importance in high-temperature applications such as jet aircraft engine components [106] under conditions of high temperature, stress, and exposure to halide salts [59, 107]. Simultaneous cycling of temperature and stress may result in reduced susceptibility to hot salt SCC compared to isothermal, monotonic loading exposure [106]. The mechanism of hot salt cracking resembles that of SCC in aqueous halide solutions, with a fracture process that is associated with hydrogen embrittlement [108]. F2. Corrosion Fatigue Titanium’s superior corrosion resistance renders it an attractive structural material for use in many corrosive environments. In the presence of cyclic loading, the environmental effects on the fatigue properties become important. The excellent corrosion resistance of titanium in seawater and many other aqueous chloride media leads to smooth and notched fatigue run-out stresses, which are virtually unaffected by the environment [8]. Figure 61.8 illustrates the effect of seawater on the fatigue crack propagation rate. Effects of cycle frequency, stress ratio, microstructure, and applied potential have been reported in the literature [8, 109, 111, 112].

The fatigue crack growth rate was found to be essentially independent of cycle frequency in air and in noncorrosive aqueous sodium sulfate solutions [111]. In aqueous halide solutions, a frequency-related crossover effect can occur at a stress intensity range DK, or DKSCC, associated with cyclic SCC [111]. Below DKSCC, lower frequencies permit more time for repassivation of fresh metal surface at the crack tip, thereby lowering the crack growth rate. Above DKSCC, lower frequencies allow more time for hydrogen diffusion and embrittlement, thereby increasing the crack growth rate. In a 3.5% NaCl solution, a crossover effect has been observed with Ti 6 Al–4 V but not with Ti 8 Al–1 Mo–1 V, indicating a relation to alloy chemistry [112]. In methanolic halide solutions, fatigue crack growth rates have been found to increase with decreasing frequency over the whole range of DK [111]. In air and saltwater environments, a significant improvement in fatigue crack growth rates in a–b alloys has been associated with a transformed beta microstructure versus an equiaxed microstructure [112]. Under ripple load conditions, in which a small cyclic load is superposed on to a sustained load, an equiaxed microstructure exhibits better cracking resistance [113]. In air, the fracture surface appearance is predominantly ductile. In alcoholic and aqueous environments, an increasing fracture surface roughness appears linked to higher fatigue crack growth rates with cleavage fracture dominating over ductile fatigue striations [111]. Commercially pure titanium and its weld metal have displayed increasing crack growth rates with increasing stress ratio in air and natural seawater [109]. This has been attributed to crack closure effects. At a relatively high DK level, an acceleration of the crack growth rate in seawater was more pronounced for weld metal than for base metal indicating a microstructural influence. A small applied cathodic potential reduced the environmental effect by suppressing anodic dissolution when the passive layer failed [109]. F3. Hydrogen-Induced Cracking

FIGURE 61.8. Trends of fatigue crack propagation rates for commercially pure titanium [109] and the intermediate strength near- a alloy Ti 5111 [110] in air and seawater.

The oxide film on titanium is an excellent barrier to hydrogen gas intrusion. Disruption of the oxide film allows easy absorption of hydrogen in high-pressure/temperature anhydrous gas streams. However, small quantities (2%) of moisture or oxygen immediately passivate the surface, forming again an effective barrier at temperatures as high as 315 C and pressures up to 800 psi [114]. At highly acidic or alkaline conditions, corrosion processes affect the integrity of the surface oxide and the corrosion potential drops below the hydrogen evolution potential. Such a drop in the open-circuit potential may also be the result of cathodic protection [115, 116], galvanic coupling, or intense dynamic abrasion. Electrochemically produced atomic hydrogen can now be absorbed. At temperatures 9000 ppm) [119]. As a result of such vast differences in hydrogen solubility in the different a and b phases, strong effects of the microstructure on the hydrogen-induced crack growth susceptibility are observed in near a and a–b alloys. Alloys with a continuous b phase provide a fast diffusion path for hydrogen, and these alloys are more susceptible to hydrogen embrittlement than alloys

874


with a continuous a phase [126]. Similar to slow strain rate embrittlement, a sustained load will promote crack growth in Ti 6-4 with increasing hydrogen levels. However, Ti 6-4 tensile properties are not affected by hydrogen levels up to 300–600 ppm [119]. The high solubility of hydrogen in the b phase renders b alloys rather insensitive to hydrogen embrittlement. Whereas several thousand parts per million may be required for any significant loss of ductility, the hydrogen absorption rate is also much higher, resulting from the much hydrogen diffusion coefficient in the b phase [127]. Hydriding can be avoided if proper consideration is given to equipment design and service conditions. At temperatures ~80 C, detrimental galvanic couples should be eliminated and impressed cathodic potentials below the hydrogen evolution line should be avoided. At all temperatures, a pickled surface provides a greater resistance to hydrogen uptake than a sandblasted, abraded, or otherwise damaged surface [114]. Disruption of the surface oxide by smeared iron permits entry of hydrogen at any pH level. In the presence of CO2, carbonate films reduce hydrogen absorption of titanium, illustrating the importance of the nature of the surface film. Small concentrations of H2S had no effect [128]. Alloying elements, such as 0.8% Ni in Grade 12, can account for increased hydrogen absorption [128]. At temperatures 6.5

pH > 6

pH > 6 pH > 6.5

Neoprene

Soft Natural Rubber

NR NR NR NR 200 200 200 150 200 200 150 200 NR NR

NR NR NR NR 150 150 150 150 125 150 150 150 NR NR

NR NR

150 180 NR NR 200 150 200 ID NR NR NR 125

pH > 6

150 175 200

Butyl

Hypalon

NR NR NR NR 200 200 200 200

150

NR NR

NR 150

NR 175

200 200

175

200 200 NR NR

185 185 200 150 NR

150 150 175

NR

NR

150

NR

100

150

ID

100

100 150 150 NR NR 100 125 150 100 ID 100 NR NR NR 150

175 185 200 NR NR 150 200 200 150 ID 150 NR NR NR 150 150

NR NR NR 150

100 150 150 NR 100 150 150 NR NR 150

NR 200 200 200 NR NR NR

NR 150 150 150 NR NR NR

pH > 6.5 200

Semihard and Hard Rubber

200 150 200 175 200 NR NR 150 NR 200 200 200 NR NR NR

185 200 150 150 200 150 185 ID 150 NR ID NR 150

185 ID

NR

175 ID NR 175 175 ID NR ID NR

200 200 200 200

NR NR

NR NR

100 185 200 NR NR NR

200 200 200 NR NR NR (continued )

938

USING PLASTICS, ELASTOMERS, AND COMPOSITES FOR CORROSION CONTROL

TABLE 66.12. (Continued )

Chemicals Butyl alcohol (butanol) Butyric acid Cadmium cyanide Calcium acetate Calcium bisulfate Calcium bisulfite Calcium bleach Calcium carbonate Calcium chloride Calcium hydroxide Calcium hypochloride Calcium nitrate Calcium oxide, dry Calcium sulfate Carbolic acid (phenol) Carbon bisulfide Carbon dioxide (wet) Carbon dioxide (dry) Carbon tetrachloride Carbonic acid Castor oil Caustic soda Chloracetic acid Chlorinated hydrocarbons Chlorine, dry Chlorine, wet Chlorine dioxide Chromic acid Citric acid Copper carbonate Copper chloride Copper cyanide Copper nitrate Copper sulfate Cottonseed oil Cresylic acid Ethanol Ethers Ethyl acetate Ethyl alcohol Ethyl chloride Ethylene glycol Fatty acids Ferric chloride Ferric hydroxide Ferric nitrate Ferric sulfate Ferrous ammonium sulfate Ferrous chloride Ferrous hydroxide Ferrous nitrate Ferrous sulfate Fluoboric acid

Remarks

Neoprene NR NR 150 150 NR

(Calcium hypochlorite) pH > 6 pH > 6.5

(Sodium hydroxide)

pH > 6.5

(Ethanol)a

(Ethanol)a

pH > 6

200 175 200 NR 200 200 150 NR NR 200 200 NR 200 ID 200 NR NR NR NR NR NR 150

150 200 NR NR 100 NR NR 100 NR 100 NR 85

pH > 6.5 200 pH > 6

150 100

Soft Natural Rubber


100 NR 100 NR

Butyl

Hypalon

150 100 150 NR

ID NR

NR NR 150 150 150

100 NR

150 NR

120 125

125

100 150 150 NR 150 150 150 NR NR 150 150 NR 150 ID 150 NR NR NR NR NR NR ID

200 200 175 NR 200 200 200 NR NR 175 175 NR 200

185

200

125 185

200

125 NR 150 NR NR 100 NR NR 100 NR 100 NR 150 100 NR 150 150 100 100 100 150

200 100 NR SC SC NR NR 150 200 100 150 150 150 NR NR 100 NR NR 100 NR 200 NR 200 150 150 200 200 200 150 200 200

100 NR 175 175 NR 185 150 185 120 NR NR NR NR ID

175 NR NR 200 200 NR 200 200

NR NR NR 125 100

185 185 185 100 NR 100

100 NR 185 NR 185

185 185

185

150 200 NR NR ID NR NR ID NR 150 NR 100 125 125 200 175

175 125

BARRIER APPLICATIONS (LININGS AND COATINGS)

939


Chemicals Fluorine gas (wet) Fluorine gas (dry) Fluosilicic acid Formaldehyde, 5% Formaldehyde, 40% Formic acid Gasoline Glauber’s salts Glycerine Hydrobromic acid Hydrochloric acid Hydrofluoric acid Hydrofluosilisic acid Hydrogen peroxide Hydrogen sulfide Hydrogen sulfite, dry Hydrogen sulfite, wet “Hypo” photographic solution Hypochlorous acid Kerosene Lacquer solvents Lactic acid Lead chloride Lead nitrate Lead sulfate Lime, dry Lime, slaked Linseed oil Lithium chloride Lye Magnesium carbonate Magnesium chloride Magnesium citrate Magnesium hydroxide Magnesium nitrate Magnesium sulfate Maleic acid Malic acid Manganese ammonium sulfate Manganese chloride Manganese sulfate Mercuric chloride Mercuric cyanide Mercuric nitrate Mercurous nitrate Methyl alcohol Methyl chloride Mineral oils Muriatic acid Nickel acetate Nickel ammonium sulfate Nickel chloride Nickel nitrate Nickel sulfate

Remarks

(Sodium sulfate)

(Sodium thiosulfate)

Pure pH > 6 pH > 6.5 (Calcium oxide) (Calcium hydroxide) pH > 6 (Sodium hydroxide) (Basic) pH > 6 pH > 6.5

pH > 6

(Methanol)a

(Hydrochloric acid) pH > 6 pH > 6 pH > 6.5

Neoprene NR NR 100 NR NR NR NR 200 NR NR NR 100 NR NR NR NR 100 NR NR NR 80 200 100 200 200 NR 200 200 200 200

Soft Natural Rubber


Butyl

Hypalon

NR NR 185

NR NR 125

125 150 NR 150 150 150 NR 150 185 NR NR

NR NR NR 200

NR NR NR 150

NR NR NR ID 150

NR NR 150 NR NR NR NR 150 100 100 150 NR 100 NR NR NR NR 150 NR NR NR 120 100 100 100 150 150 NR

NR NR 200 150 150 NR NR 200

150 200 175 NR

180 150

150 100 150

200

158 185

NR 150 175

185 185 185

150 150 175

200 200 NR 200 NR NR NR NR 200 NR NR NR 150 150

150 180

NR NR NR 125 NR NR

175

200 200 200

150 150 150

200 150 200 200 200

ID

NR

150

NR

80

200 NR NR

150 100 100

NR ID NR NR NR ID

NR 80 NR NR 150 100

150 200 150 150 150 150 NR NR 180 200

180 150 150 ID ID 180 NR NR NR 150

175 ID ID ID ID 80 NR NR NR 125

200 100 200

150 100 150

200 200 200

150 ID 185

175 150 200 (continued )

940



Chemicals

Remarks

Niter

(Potassium nitrate) pH > 6.5

Nitric acid, 5% Nitric acid, 10% Nitric acid, 25% Nitric acid, 40% Nitrous acid Oleic acid Oleum Oxalic acid Palmitric acid Paraffin Perchloric acid Peroxide bleach Petroleum oils, crude Phenol Phosphoric acid, 85% Plating solution, brass Plating solution, cadmium Plating solution, chrome Plating solution, copper Plating solution, gold Plating solution, lead Plating solution, nickel Plating splution, silver Plating solution, tin Plating solution, zinc Potassium acetate Potassium acid sulfate Potassium alum Potassium aluminum sulfate Potassium antimonite Potassium auricyanide Potassium bicarbonate Potassium bichromate Potassium bisulfate Potassium bisulfite Potassium borate Potassium bromide Potassium carbonate Potassium chlorate Potassium chloride Potassium chromate Potassium cyanide Potassium cyprocyanide Potassium dichromate Potassium ferricyanide Potassium hydroxide, 25% Potassium hydroxide saturated Potassium iodide Potassium nitrate Potassium nitrite Potassium oleate Potassium permanganate

(Fuming sulfuric acid)

(Dihydrate) (Sodium perborate) (Carbolic acid) (Over 85% use butyl)

Neoprene

Soft Natural Rubber


200

150

180

ID

200

NR NR NR NR NR NR NR 150 NR

NR NR NR NR NR NR

NR NR NR NR 150

150 125 100 NR

100 NR NR NR NR

NR

NR

NR NR 150

NR NR 150

NR

NR

200

150 150

200 150 150 NR

100 NR 150 NR 100 100 100

Hypalon

150

NR 175

ID

NR

NR NR 200 200 200 NR 200 200 200 200 200 200 200

NR NR 185

NR

NR

NR

180 175

200 200

150 NR 125 125 NR 100 150 150 150 150

200 200 150 150 180 NR 200 200 200 200 200 200 200 NR 200 200 NR NR 200 150 200 180

NR

150

NR NR

NR 150 NR NR NR

Butyl

pH > 6

(Alum)

pH > 6

200 pH > 6 pH > 6

150 150

pH > 6 pH > 25% pH > 6.5 pH > 6.5 pH > 6.5 pH > 7.0

NR NR 220 220 200

175

175 ID

180 NR 150 NR ID 185

180

200 200 200 NR 150 175 ID NR NR 150 200

NR

BARRIER APPLICATIONS (LININGS AND COATINGS)

941


Chemicals

Remarks

Potassium phosphate

Mono-di- or tribasic

Potassium salicylate Potassium silicate Potassium sulfate Potassium sulfide Potassium sulfite Potassium thiosulfate Propane Propionic acid (dilute) Propyl alcohol Rochelle salts Sludge acids Soap solutions Sodium acid sulfate Sodium aluminum sulfate Sodium antimonate Sodium bicarbonate Sodium bichromate Sodium bisulfate Sodium bisulfite Sodium borate Sodium carbonate Sodium chlorate Sodium chloride Sodium cuprocyanide Sodium cyanide Sodium dichromate Sodium ferricyanide Sodium fluoride Sodium hydroxide, 25% Sodium hydroxide, saturated Sodium hypochlorite Sodium nitrate Sodium nitrite Sodium oeate Sodium perborate Sodium permanganate Sodium peroxide Sodium phosphate Sodium salicylate Sodium silicate Sodium sulfate Sodium sulfide Sodium sulfite Sodium thiosulfate Soft drink syrups and concentrates Stannic chloride Stannous chloride Stearic acid Sulfate liquors Sulfite liquors Sulfur dioxide, wet

pH > 6

Neoprene

Soft Natural Rubber


200

150

200

200 200 200 150 AMB NR

100 150

Butyl

Hypalon 200

150 100 NR NR 100 100

200 200 200 200 200 200 NR NR 150 200

100

180

100

200 200

150

200

200

100 150

200 200

200

150

200

185

150

100 NR 100 NR 150

200 NR

150

NR 200

185

NR

NR 150

NR 180

100 180

200

150 NR NR 150

150 200 NR 200 200 200 200 200 200 200

(Potassium sodium tartrate)

105

200

NR

NR

200

200 pH > 6

pH > 6

200 150 170

pH > 6 NR pH > 25% pH > 9 pH > 6.5 pH > 6.5

200 200 NR 200 150 200

pH > 7.0 Mono-di- or tribasic

200 200 200

pH > 6

150

NR pH > 6 pH > 6

NR NR NR

150 150 100 150 150 NR 150 150 NR

200

200 200 150 ID

200 200 NR

NR 185

200

150

200

150

150

125 150

150 (continued )

942



Chemicals Sulfuric acid, 5% Sulfuric acid, 25% Sulfuric acid, 50% Sulfuric acid 75% Sulfurous acid Tannic acid Tartaric acid Tin chloride Trichloroethylene Triethanolamine Trisodium phosphate Turpentine Urea Urea ammonia liquor Urec acid Vegetable oils Water, acid mine Water, fresh Water sea or salt White liquor Wood pulp and pulp liquors Zinc acetate Zinc chloride Zinc nitrate Zinc sulfate

Remarks

Neoprene

Soft Natural Rubber


180 170 75 NR NR 1D 100

100 100 NR NR NR 100 150

NR 200 NR

Butyl

Hypalon

185 175 150 NR 200 150 200

185 175 150 NR 150

175 175 150

175

125

NR

NR

NR

NR

150 NR

NR

NR

NR

200

NR 100 150 150

150 185 200

200

100 200

150

150

200

185

150

150

130

200

185

175

Ph > 6 pH > 6

150

pH > 6

a

No harmfull impurities.

Loose linings are fabricated by making cylinders by using butt fusion machines. Axial welding of the cylinders is performed by hot air welding. Nozzles are welded either by butt fusion or hot air welding. Flaring of the nozzle is performed after the liner is inserted in the housing. Performance for loose linings has been fair in larger vessels. For obvious reasons this system is not used for vacuum service. Such linings have been used in service with fluorobenzene. 5. Molded Linings. For components like pumps and valves, melt-processible fluoropolymer linings are often formed by transfer molding. Sizes are limited by the capacity of the transfer molding equipment, polymer rheology, and other factors.

Special molding technology is used to form PTFE linings for components and piping. Although PTFE is a thermoplastic, it is so viscous in the melt that conventional molding techniques cannot be used. To line components, the resin is pressed in housings, usually by isostatic pressure applied by inflating rubber tooling. The resulting “green form” is sintered at high temperature so that it coalesces into a solid shape. Isostatic molding is also used to form PTFE shapes for pumps and other components. In most cases, these shapes require machining to meet required tolerances. Some PTFE pipe linings are formed by a paste extrusion process. The resin is mixed with a hydrocarbon extrusion aid and extruded as tubing. The extruded shape is sintered to drive off the hydrocarbon and coalesce

TABLE 66.13. General-Purpose, Thick Thermoplastic Sheet Liningsa Material

Forum

Description

Polypropylene Plasticized PVC Unplasticized PVC on plasticized PVC

Rubber-backed sheet Sheet Laminated sheet

0.060-in.-thick Polypro on 0.060 in. rubber 1 3 3 16 in., 32 in and 16 in. thick 0.085 in. thick

a

Root pass welding followed by capstrip covers is used.

TABLE 66.14.

Fluoropolymer Physical Propertiesa

Property

Test Method

Units

PTFE

FEP

PFA

ETFE

ECTFE

PVDF

Tensile strength Elongation Flexural modulus Hardness Melting point Upper service temp.

D638 D638 D790 D2240 DTA, E-168 UL746B

psi 1000 (MPa) % psi 1000 (MPa) Shore D F ( C) F ( C)

3–5 (21–35) 300–500 72 (500) 50–65 621 (327) 500 (260)

3.4 (23) 325 85 (600) 56 500 (260) 400 (204)

3.6 (25) 300 85 (600) 60 582 (305) 500 (260)

5.8–6.7(40–47) 150–300 170 (1200) 63–72 473–512 (245–267) 300 (150)

4.2–4.3 (28.9–29.6) 200 240–245(1655–1689) 75 464 (240) 300 (150)

6.8–8.0(46.9–55.2) 50–250 165–325(1138–2241) 77 320–338 (160–170) 275 (135)

Arkema sells PVDF as Kynar. a Sources: DuPont (PTFE, FEP, PFA, ETFE) and Solvay Solexis (ECTFE, PVDF). DuPont sells PTFE, FEP, and PFA as Teflon fluoropolymer resins and ETFE as Tefzel fluoropolymer resin. Solvay Solexis sells ECTFE as Halar and PVDF as Hylar.

943

944


TABLE 66.15. Chemical Resistance of Partially Fluorinated Fluoropolymer Manufacture Recommended Maximum Service Temperature

(continued )

the resin. Other processes for PTFE are ram extrusion and tape-wrap, where layers of tape are built up on a mandrel to the required thickness before sintering. 6. Dual Laminates. Structures are made of FRP built up on fluoropolymer sheeting that serves as a lining. Dual-laminate structures combine the technology of fluoropolymer fabrication with the technology of FRP (see Section G for self-supporting structures). This lining is fabricated by machine and hand welding of fabric-backed fluoropolymer sheets like those described above for adhesive-bonded and loose

linings. The fabric aids in bonding the sheeting to the FRP. Dual laminate construction involves fabricating the thermoplastic shell using the welding techniques described earlier for adhesive bonds and loose linings. The FRP shell is then wound on the thermoplastic structure followed by welding of the nozzles and an FRP overlay on the joints. Some commonly used liners are PVDF, ECTFE, ETFE, FEP, and PFA.

SELF-SUPPORTING STRUCTURES: PROCESS VESSELS, COLUMNS, AND PIPING

945


(continued )

Table 66.17 compares the five technologies of Fluoropolymer lining.

G. SELF-SUPPORTING STRUCTURES: PROCESS VESSELS, COLUMNS, AND PIPING G1. Thermoplastics The use of thermoplastic polymers in self-supporting structures is limited to piping. The exceptions are rotomolded

polyethylene and welded polypropylene tubs. Materials used for piping are PVC, CPVC, PE, and PP. On rare occasions PVDF piping is also encountered. These materials are extruded for piping that is available up to 356 mm (14 in.) diameter. G1.1. Design of Piping. Thermoplastic piping is designed on the basis of the hydrostatic design basis (HDB), which involves carrying out short-term burst tests at various pressures and recording the time to failure. The stress versus time to failure is extrapolated to an arbitrarily chosen long-term

946



(continued )

value called the HDB. A factor of safety is applied to calculate the maximum allowable stress and pressure. G1.2. Joining of Pipe. Joining of piping products such as lengths and fittings is done by welding (for PE, PP, and PVDF) and solvent cementing for PVC. Solvent converting is a quick and easy way to join pipe lengths and fittings through couplings. Welding requires qualified personnel using standardized industry practices.

G1.3. Welding of Thermoplastics. There are three principal methods of welding thermoplastic materials. Welding is needed to join piping ends on edges of sheet such as in lining: Hot air (hand) welding Extrusion welding Butt fusion (hot plate or heated tool) welding Flow fusion welding Cap strip welding


947


(continued )

Hot air and extrusion welding are similar except that extrusion weld is used for sheets thicker than 1/8 in. In both cases a weld rod is fused between the two sheets to be joined. In the case of hot air welding the weld rod is manually pressed in a V-groove while hot air or nitrogen melts the skin of the sheet and the rod. Sheet linings for fluoropolymers, adhesively bonded, PP and PVC use this technique. In extrusion welding the filler rod is extruded in the groove by an extruding screw.

Butt fusion welding is most used for welding pipe ends of both single- and double-wall piping. A special machine capable of properly aligning the pipe ends and melting the ends and pressing them together is used. This technique also used for axial welds of sheets for linings. How fusion welding is a variation of the butt fusion welding where a special machine is used which melts and fuses the edges of sheets together. This technique is used for making longitudinal welds in loose fluoropolymer linings as well as dual-laminate liners.

948



(continued )

Cap strip welding is used to cover hand welds in the bonded linings (fluoropolymers, PVC, and PP) as well as dual laminate. It is similar to hand (hot gas) welding except a strip is used instead of a rod. Regardless of the welding technique used, successful welds have the following in common: controlled heat input, consistent and right level of pressure, and the time for which heat and pressure are maintained. Figure 66.11

shows the schematic of a welding gun used for any free hand welding including extrusion welding. Figure 66.12 shows how free hand welding is done using a gun and the weld rod. Fluoropolymers are welded with an electric weld gun using a speed welding tip. (Fig. 66.13). Figure 66.14 shows root pass welding followed by cap strip welding. Currently the industry follows AWS (American Welding Society) and DVS (German Welding Institute) standards


949


for weld quality and qualification. Following are some of the standards used: Weld qualification DVS certification program per DVS 2212 AWS B 2.4 (2006)

G2. Thermosetting Materials: FRP (RTP)

Weld evaluation DVS 2203 Part 5 (bend test) DVS 2203 part 4 (tensile creep test) AWS G 1.10M (2001)—visual examination EN 12814-7, Tensile test with waisted test specimen G1.4. Welded Vessels. Using the hand (hot gas) welding techniques PP tubs are fabricated and are widely used for TABLE 66.16.

Permeation Rules of Thumb Change

Voids in polymer Polymer crystallinity Polymer chain stiffness Polymer interchain forces Polymer temperature Permeant size/shape Permeant concentration Polymer thickness Permeant temperature Voids in polymer Permeant/polymer chemistry

+ * * * + * + * + * + Similarity

pickling operations. By far the most commonly used vessels are rotomolded PE (HDPE as well as crosslinked) per ASTM D 1198. These are economical and find use in HCl and H2O2 storage.

Permeation + + + + + + + + + * +

The bulk of self-supporting structures are made of reinforced thermosetting materials commonly known as fiberreinforced plastics (FRPs or reinforced thermosetting plastics (RTPs.) Resin constitutes the continuous phase (matrix) and fibers (continuous or chopped strand as well as mat or woven roving) form the reinforcement. Resin provides the needed chemical resistance while the reinforcement provides the needed mechanical and physical properties such as strength, elongation, and insulation. G2.1. Resins. The resins used are generally known as polyesters. In reality, resins belong to the following broad classifications: Bisphenol A fumarate Chlorendic anhydride Epoxy vinyl esters Epoxy Novalac–based vinyl esters Modified vinyl esters Furan Table 66.18 shows the building blocks of these resins.

950 TABLE 66.17.

Fluoropolymer Lining Systemsa

Lining System

Lining Materials

Maximum Size

Design Limits

Fabrication

Adhesive bonding

Fabric-backed PVDF, PTFE, FEP, ECTFE, PEA

No limit

Neoprene or epoxy adhesive, sheets welded. Heads thermoformed or welded

Dual laminate

Same as adhesive bonding

12 ft (3.7 m) dia.

Sprayed dispersion

FEP, PFA, PFA w/mesh and carbon, PVDF, PVDF w/glass or carbon fabric ETFE, FEP, PFA, ECTFE, PVDF

8 ft (2.4 m) dia., 40 ft (12.2 m) length

Pressure allowed. Full vacuum only at ambient temperature. Smallest nozzle is 2 in. (51 mm). Maximum temperature limited by adhesive, typically 275 F (135 C) No pressure allowed. Vacuum rating not determined Pressure allowed. Vacuum rating not determined

8 ft (2.4 m) dia., 40 ft (12 .2 m) length

Pressure allowed. Vacuum rating not determined

8 ft (2.4 m) dia., 22 ft (6.7 m) length

Pressure allowed. Vacuum rating not determined Pressure allowed. Vacuum rating depends on lining thickness Pressure allowed. No vacuum. Gasketing required between liner and flange face

Electrostatic spray—powder Rotolining

ETFE, PVDF, ECTFE

Isostatic molding, paste extrusion Loose lining

PTFE

a

FEP, PFA

Determined by body flange

Linear fabricated on mandrel. FRP built up over liner Multicoat application. Each coat is baked Multiple coats applied by electrostatic spraying. Each coat is baked Rotationally molded PTFE performed under pressure or paste extruded, then sintered Liner with nozzles hand or machine welded

Nondestructive spark testing should be used, along with visual inspection for all systems except loose linings. Adhesive bonding can be done in the shop or field; other systems are shop only.


951

FIGURE 66.11. Electrically heated welding gun.

Styrene is added to these resins as a reactive diluent. Styrene serves two purposes: First, it helps in adjusting the viscosity of the resin mix for ease of fabrication and, second, it lowers the cost of the resin. The final structure is a copolymer of styrene and the polyester resin. Choice of resin will depend on chemical resistance, temperature resistance, flexibility, and ease of fabrication. Resin manufacturers routinely publish chemical resistance information as well as mechanical properties. Bis A fumarate resins offer the best resistance to caustic solutions but lack the flexibility of modern resins. These resins were the early workhorse in CPI and have been virtually replaced by the more versatile vinyl ester resins which offer a variety of combinations of chemical resistance, flexibility, hightemperature resistance, and ease of fabrication. The

chlorendic anhydride family is principally used for chlorine-handling applications while furans are used for solvent resistance. Currently, vinyl ester resins account for the largest number of applications in the CPI. G2.2. Reinforcements. The FRP laminates are heterogeneous (i.e., the type and the amount of reinforcement varies across the thickness.) The surface against the chemical is reinforced by a glass or synthetic veil (polyester or carbon) and is resin rich (90% resin, 10% glass). Next are two or four layers of chopped strand mats with a somewhat lower resin/ glass ratio (70% resin). Together, the veil(s) and mats form the corrosion barrier. The structural thickness that follows is either by alternating layers of mat and woven roving known as contact molding (hand layup) or continuous strand roving

FIGURE 66.12. Welding positions.

952


FIGURE 66.13. Electric gun with speed welding tip.

interspersed with uniaxial fabric and spray up (filament wound construction). [see Figs. 66.15(a) and (b)]. G2.3. Other Additives. For fire retardancy, resins are brominated. Antimony trioxide is also added. G2.4. Curing Systems. There are three principal curing systems used for FRP. Methyl ethyl ketone peroxide (MEKP) is by far the most common. For higher crosslink density and the consequent increase in chemical resistance, benzoyl peroxide (BPO) with dimethylaniline (DMA) is used. For thicker sections, where a slower cure is desired, cumine hydro

peroxide (CHP) is used. Curing refers to the resin crosslinking process caused by free radicals formed by the curing agents. Sometimes postcuring is required to increase the degree of crosslinking which enhances the chemical resistance. Hardness measurements by the Barcol tester or the acetone sensitivity test measure the degree of crosslinking and the chemical resistance. G2.5. Selection of Resin and Laminate Composition. Selection of resin is currently done by the ASTM C581 test. It involves plotting the retained flexural strength of coupons against time. The resin is accepted if the curve levels off at no more than 50% reduction. Currently, efforts are underway to develop enhanced resin selection techniques. Quite often, the selection of resin and the laminate design is determined by successful case histories in the industry or actual field tests. Table. 66.19 indicates general chemical resistance of FRP resins. G2.6. Materials for Dual-Laminate Construction. Dual laminate refers to FRP structure with a bonded thermoplastic liner on the inside. The liner materials range from low-end PVC through PP, PVDF (Kynar ), ECTFE (Halar ), ETFE (Tefzel ), FEP (Teflon ), and PFA. When using the low-end thermoplastics such as PVC and chlorinated polyvinyl chloride, it is possible to use low-grade isophthalic polyester resins. This practice is quite common in Europe and to a lesser extent in Canada. In the United States, the cost of standard vinyl ester resin is considered to be insignificant in proportion to the total cost of the dual-laminate structure. Additionally, the higher grade resin provides some protection in the event of failure of the thermoplastic liner.

FIGURE 66.14. Root pass and cap strip welding for fluoropolymer lining.

G2.7. Selection of Thermoplastic Liners. Selection of the thermoplastic depends on the chemical and temperature resistance required, weldability of the thermoplastic, and cost. Dual laminates must be considered in conjunction with


953

TABLE 66.18. Building Blocks for FRP Resins Generic

Acid

Glycol

Unsaturated Acid

Trade Namesa

ISO

Orthophthalic Isophthalic Terephthalic

Fumaric Maleic anhydride

Atlac 400

Bis A furmarate

Bisphenol A

Fumaric acid

Atlac 382

Vinyl ester

Bisphenol A

Ethylene Propylene Diethylene Dipropylene Neopentyl Ethylene oxide Propylene oxide Epichlorohydrinb

Methacrylic acid

Epoxy Novoloc Vinyl

Phenol formaldehyde

Epichlorohydrinb

Methacrylic acid

Ester Chlorendic anhydride

Chlorendic acid

Ethylene oxide Propylene oxide

Maleic anhydride

Atlac 580 Derakane 411 Hetron 922 Derakane 470 Hetron 970 Hetron 92

a b

Atlac is a registered trade name of Reichhold Co. Hetron and Derakane are registered trade names of Ashland Chemical Co. Not a glycol.

other alternatives such as metals and ceramics. Manufacturers of thermoplastics publish chemical resistance data. In the event that testing is needed, dunk tests of thermoplastic materials are carried out for weight change and retained mechanical properties. More appropriately laboratory (ASTM C868) or field blind flange tests should be carried out. Loss of adhesion from FRP, color change, embrittlement, and stability of welds should be examined. G2.8. Design and Fabrication of FRP Vessels and Columns. The FRP structures are designed on the basis of operating condition temperatures and pressures (imposed as

well as hydrostatic) as well as other loads such as seismic, wind, snow, attached piping, and agitator. Design is carried out by established industrial practices and standards such as ASTM D3299, D4097, RTP-1, and ASME Section X. Complete design includes shell and head thicknesses, type of heads, joints for shell sections, nozzles to shell, and shell to heads. Design of lifting, support lugs, and anchors is also included. The FRP vessel fabrication is performed by one of the following two techniques. G2.8.1. Contact Molding (Hand Layup). This technique involves sequential layup of layers saturated with resin on

FIGURE 66.15. Chemical resistant structural laminate: (a) Filament wound; (b) hand lay up.

954

TABLE 66.19. Chemical Resistance of FRP Bisphenol A Epoxy, Amine Cured

Bisphenol A Epoxy, Anhydrous Cured

Bisphenol A Vinyl Ester

Novolac VE

Bisphenol A Fumarate Polyester

Chlorendic Acid Polyester

Furan

Isophthalic Acid Polyester

Orthophthalic Acid Polyester

Up to Temperature ( F)

120

210

120

210

120

210

120

210

120

210

120

210

120

210

120

180

120

150

Acetaldehyde Acetaldehyde, aq. 40% Acetic acid, glacial Acetic acid, 20% (25) Acetic acid, 80% Acetic anhydride Acetone, 10% Adipic acid Alocohol, ally1 Alcohol, benzy1 Alcohol, buty1 (n-butanol) Alcohol, buty1 (2-butanol) Alcohol, ethy1 Alcohol, hexy1 Alcohol, isopropy1 (2-propanol Alcohol, methy1 Alcohol, propy1 (1-propanol) Allyl chloride Alum Ammonia, gas Ammonia, liquid Ammonia, aq. 20% Ammonium salts, except fluoride, Ammonium fluoride, 25% Amyl acetate Amyl chloride Aniline Aniline hydrochloride Antimony trichloride Aqua regia Arsenic acid, 80% Aryl-sulfonic acid

N N C R C C N C N C C

N N N C N N N N N N N

N N N R N N N C N N C

N N N C N N N N N N N

N N N R C N N R N N C

N N N R N N N N N N N

R N C R R C R R N C R

N N N R R N N N N N N

N N N R N N N

N N N R N N N

N N C R N C N

N N N R N N N

R R N C N R N

R R N C N R N

N N N R N N N

N N N N N N N

N N N N N N N

N N N N N N N

N N N

N N N

N N N

N N N

R R

N N N

N N N

N N N

N N N

C

N

N

N

C

N

R

N

N

N

N

N

R

N

N

N

N

R R R

C C N

C R C

N C N

C C R

N N N

C R R

N C N

C N N

N N N

N N N

N N N

R R R

R N N

N N N

N N N

N N N

C R

N N

N R

N N

N C

N N

C R

N N

N N

N N

N N

N N

R R

R

C N

N N

N N

N N

N R C N

N C N N

N R C N

N C N N

N R N N C R

N R N N N R

N R N N N R

N R N N N C

R

C

N R N N N R

R

R

N R C N R R

R

C

N R N N N R

R

R

N R C N R R

R

R

N R R N N R

N R N N N R

N N N N N R

N N N N N R

R

N

R

N

R

N

R

N

R

N

R

N

N

N

N

N

N

N

C R N R

N N N N

N R N R

N N N N

N N N N

N N N N R N

N R

C R

N R

N N N N N N N R

N N N N C N

C R

N N N R N R N R

N N N N R R

N R

N N N N C N N R

R R N R

C R

N N N R R N C R

N

N

N N N N N N N N

N N N N N N N N

N N N N N N N N

R R R

R N

R R R

N

955

Barium salts Beer Beet sugar liquor Benzaldehyde, 10% Benzaldehyde, 10–100% Benzene (benzoil) Benzenesulfonic acid,10% Benzenesulfonic acid,50% Benzoic acid Black liquor––paper Bleach, 12.5% active chlorine Bleach, 5.5% active chlorine Borax Boric acid Brine Bromic acid, 20 dry mils). Impermeability to ionic solutions and oxygen is an entirely more practical objective, however, and this factor is thought to be rate determining for corrosion beneath barrier films. Permeability of the film to water is generally thought to have greater direct consequence on deadhesion, leading subsequently to corrosion. C1. Binders for Barrier Coatings C1.1. Thermoplastic Binders. Impermeability is ensured primarily by the molecular structure of the binder [19, 20], although the selection of pigments and other components of the coating (volatiles as well as nonvolatiles) can be critical [21, 22]. Dense, molecularly tight films (thermosetting films of uniform high crosslink density or relatively crystalline thermoplastics) that are hydrophobic and give optimized resistance to the ingress of water, oxygen, and ionic material are the preferred binders. High-molecular-weight hydrocarbons, particularly halogenated species (fluoropolymers, polyvinyl chloride, etc.), are particularly suited to this employment. Unfortunately, because of the high crystallinity of these systems, the polymers are not soluble in conventional solvent systems, and, where they are, the low concentration of the polymer at practical solution viscosities excludes their use in many applications where restrictions apply to the amount of solvent which may be employed. Their use in the form of dispersions from water and as powder coatings has been widely practiced in coil coatings, container coatings, and so on, although these systems cannot readily be employed in field applications because of thermal curing requirements. Such is the cohesive strength of many of these polymers (e.g., polyvinyl chloride) that adhesion under conditions of

field-applied stress may be tentative unless the polymer is suitably modified. At some compromise with performance, chemically engineered modifications to the polymer (plasticization, copolymerization, introduction of carboxylic acid groups, etc.) and reductions in molecular weight improve both solubility and adhesion. These modifications enable coatings based on such polymers to be applied as lacquers and dried in the field by evaporation of the solvent carrier. Nevertheless, these systems often fail to satisfy modern environmental requirements that limit the amount of volatile solvents in the coating. Additional difficulties with chlorinated polymers are related to the propensity of such materials to undergo heat- and light-induced dehydrochlorination, producing HCl and chlorides at the substrate that effectively short circuit the barrier properties of the film. Finally, solvent retention in applied thermoplastic lacquers may adversely affect the impermeability of the film to water, where retained solvents are hydrophilic. Such hydrophilic solvent entrapment will reduce the glass transition temperature Tg and may induce osmotic blistering failure. C1.2. Thermosetting Binders. In consequence of all these things, with the possible exception of high build bituminous cutbacks, most modern barrier systems are based on thermosetting polymers (epoxies, polyesters, vinyl esters, and polyurethanes) [4]. These take the form of two-pack systems that are polymerized in situ from low-molecular-weight oligomeric and monomeric precursors or single-pack thermosets (epoxy/phenolics, amino formaldehyde resin crosslinked acrylics, etc.) applied as premixed systems and polymerized by baking after application. Single-pack, thermally cured, thermosetting powder coatings may also be used, where practical, although the corrosion resistance

974

CORROSION CONTROL OF STEEL BY ORGANIC COATINGS

properties of these systems are often poorer than their twopack, solvent-based counterparts. All of these systems allow the deposition of higher solids, environmentally more desirable films of much higher thickness than is possible with the thermoplastics. With many of these systems, simultaneous control over cure kinetics and solvent release allows for enhanced adhesion and impermeability as does polymer structure (the presence of polar groups) and the mechanism of cure, that is, polymerization after (but not before) application. Some systems (vinyl esters, unsaturated polyesters), as well as those binders crosslinking via condensation-type polymerization mechanisms (phenolics and other formaldehyde-based systems), exhibit high internal stress from shrinkage on polymerization (especially when applied at excessive film thicknesses) and this may be disadvantageous to sustained adhesion [23, 24]. Epoxies, however, curing by addition polymerization, shrink less and exhibit much improved adhesion, making very valuable binders for barrier coatings. Some thermosets (epoxies and polyurethanes) are modified with crude, highly hydrophobic, bituminous-based thermoplastics (coal tar resins) and give adhesive and highly impermeable multicoat systems for optimum performance under immersion and below grade service [4, 25]. With all of these thermosets, as with the prepolymerized thermoplastics, design objectives are to create a dense hydrophobic molecular matrix through which the transport of water and oxygen is minimized and that of ionic species is entirely prevented [26, 27]. These requirements are at best satisfied by polymeric films having high uniform crosslink density, high Tg, and molecular structures that are either nonpolar or derived from non-water-attracting moieties [28].

Mayne and coworkers [29, 30] showed that in thin barrier films ionic ingress to the substrate occurs at areas of low crosslink density (in otherwise high-crosslink-density films) and it is at these sites of reduced crosslink density that corrosion is initiated. In film areas having high crosslink density, the film picks up only water. C2. Pigments for Barrier Films In practical paint films, it may be anticipated that the transfer of both water and ionic solutions into the film is also favored by the presence of other hydrophilic components (pigments and additives with high water solubility, solvents with high water miscibility). These materials should be excluded from the formulation. Pigments, as well as other components (additives, etc.), should be selected so as to enhance hydrophobicity. Wherever possible, adhesion across pigmentary/ binder interfaces should also be maximized and pigment flocculation minimized. Impermeability is enhanced by the use of flat platey pigmentations, especially aluminum flake, glass flake, stainless steel flake, and micaceous iron oxide. In wet films, these pigmentations orient themselves parallel to the substrate and both laterally reinforce the film and present a more tortuous path for the penetration of any corrosive agent through the film [20]. Pigmentation levels for barrier coatings are less critical than are pigmentation levels for zinc and inhibitive based primers, being normally much lower (see Fig. 67.2). Excessive pigmentation levels should be avoided, especially where the metallic pigmentation is cathodic to the metal substrate. This maintains the electrical resistance of the film and prevents pitting at uncoated pinholes, which act as anodes.

FIGURE 67.2. Pigment volume criteria for classes of anticorrosive coatings.

BARRIER COATINGS

C3. Solvent Systems for Barrier Films Solvent systems should be designed so that lateral orientation of the pigment in the film is optimized but so hydrophilic solvents evaporate and diffuse from the film rapidly, if possible, before cure is complete and certainly before the film is put into service. Slow solvents that readily wet the substrate and provide enough motility in the wet film to allow the flat, platelike, barrier pigments to settle flat against the substrate are to be preferred. However, slow-drying hydrophilic solvents, especially those having high affinity for the binder or those which have nonplanar complex structures, tend to be held within the drying film for long periods. They may, in service, attract water into the film osmotically, reducing the electrical resistance of the interfacial area and leading to blistering. The presence of water and/or associations of solvent and water within the film tend to plasticize and swell the film, reducing Tg and thereby facilitating the ingress of additional water. This causes further decrease in Tg. Optimally, therefore, high-boiling solvents that are last to leave the film should be hydrophobic, free of functional groups, and planar in structure. Here, some compromise with possible conflicting requirements, such as binder solubility and the need for defect-free films, may be necessary. Care should also be taken in the selection of nonvolatile additives (pigment dispersants, surfactants, thixotropes, and flow control agents) as many of these materials may be excessively hydrophilic.

C4. Wet Adhesion It has been noted that corrosion beneath a barrier film can begin only after deadhesion has taken place, and if adhesion can be maintained under wet service conditions, then protection is assured [31]. Adhesion is also optimized by polymer design. Most coatings must rely on secondary valency bonding for adhesion. A limited number of polymer types form primary valency bonds with the metal. In both cases, adhesion may be optimized when the polymeric binder is richly endowed with polar groups (hydroxyls, carboxylic acid groups, etc.) on their molecule. It is because of this that epoxies, alkyds, and polyesters show their characteristic high levels of adhesion to metals. These groups readily associate with metal oxide groups on the exposed metal surface. In designing systems for optimized impermeability and maximized adhesion, there is some dichotomy of purpose [31, 32]. As noted above, those polymers bearing polar groups that encourage adhesion (hydroxyls, carboxylic acid groups) tend to be hydrophilic and attract water into the film. Polar ester groups (e.g., in alkyds and other oxidizing systems) and, to a lesser extent, amide groups (urethanes) are also vulnerable to alkali-induced hydrolytic cleavage and may in consequence be attacked by alkalinity developed at the cathodes of corrosion cells. The effect is seen in the

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peripheral areas around corrosion spots and breaks in these sensitive films in service, where both cohesive loss (via saponification) and adhesive loss at the coating–metal interface will result. The dichotomy of need for increased adhesion and minimized water takeup and alkali resistance is resolved to some extent by the use of multicoat systems. Here adhesion becomes of primary importance in barrier primers (perhaps at some compromise to maximized impermeability) while impermeability (and the use of less polar components) becomes paramount in barrier-type finish coats [28]. Adhesion of topcoats to primers is, if recoating is done soon after priming, more readily secured than is the adhesion of the organic primer to the metal substrate. Ester group sensitivity is best resolved by avoiding polymers based on these types of binders or at least in seeking to maximize alkali resistance via the selection of polymers with suitable polymeric architecture. Not all ester-based coatings have the same degree of vulnerability to hydrolysis as do alkyds and oxidizing systems. Polyesters, vinyl esters, and ester-based polymers prepared from polyols bearing no hydrogen on the carbon atom that is beta to the hydroxyl are very often quite resistant enough for successful service in alkaline environments. If the hydroxyl group of a barrier coating binder is bound to a relatively rigid polymeric backbone, molecular mobility after cure is much reduced. Optimally, the hydroxyls in the interfacial layer of the primer would orient themselves against the metal (associating with metal oxides on the substrate) so that they are less available for hydration, and the substrate is therefore more protected. Moreover, such rigid systems are molecularly too fixed to facilitate readily the transfer of water to the substrate because of the reduced kinetic energy of the molecule [32]. The opportunity available for access of water to the substrate is therefore limited. On the other hand, during application and wetting (while the paint film is still heavily solvated) the same hydrophilic groups should display high mobility so that substrate wetting is maximized and molecular orientation with available metal oxide groups on the surface of the metal is favorably enhanced. Thus, barrier film binder design is optimized in systems having high molecular mobility in the wet stage but high immobility when the film is cured. These characteristics may be maximized where low-molecular-weight polar monomers and oligomers are applied in wet films that may soak into the substrate before converting to high Tg films of limited mobility after cure [33]. C5. Effects of Structure on Oxygen Impermeability It has been shown that polymers rich in hydroxyls have greater impermeability to oxygen than do nonpolar polymers [19, 20]. As the control of oxygen transport would

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appear more readily achievable in practical systems than the control of water transport, it may be argued that the presence of hydroxyl groups is not entirely incompatible with good corrosion control where oxygen deprivation is the desired mechanism. C6. Film Thickness One of the most effective strategies by which good barrier protection may be assured is found in the application of additional thicknesses of paint. This may be done as highbuild single-coat applications, as are often employed in protection of buried structures with bituminous cutbacks, or more effectively in multicoat systems. In the latter case, finish coats applied over primers and primer/intermediate coat combinations render additional service, that is, either enhanced impermeability (coal tar epoxies) or for other reasons (e.g., light-stable aliphatic urethanes over high-build epoxy primers for improved weathering resistance). As film thicknesses increase, so the transfer of moisture, oxygen, and ionic species to the substrate is greatly diminished, if not the absorption of such penetrants into the film. High film thicknesses may be one of the primary defenses for corrosion where engineering structures must be buried or subject to long-term immersion. In these applications, coatings are often used in tandem with cathodic protection systems, and the coatings effectively reduce the area of the exposed metal requiring cathodic protection. Thus, cathodic protection current requirements are lowered. In these systems especially, high levels of alkali resistance are necessary if adhesion of the coating is to be maintained while the coated metal remains under cathodic protection. Light-stable barrier finishes that are resistant to ultraviolet light (UV)–induced changes (degradation) in the mechanical properties of the film also prevent gloss loss, microcracking, and other age-related film defects, which lead to reduced long-term corrosion resistance of the total system. Multicoats have practical advantages over single coats in that the superimposition of several films does much to eliminate the possibility that pinholes, holidays, and misapplications of single coats may allow direct access of the environment to the metal. It is statistically unlikely that holes in one coat will coincide with similar faults in the next, although pinholes can telegraph through two coats. Perhaps, more importantly, certainly with thermosetting (chemically curing systems) internal stress development on drying, arising from solvent release and/or from polymerization, will increase as film thickness increases [24, 34]. The stress produced by the conversion of thick films (even strongly adherent epoxies) is capable of overcoming the adhesion of the coating, even over abrasive blasted surfaces. In less extreme cases, residual internal strain left in a coating after curing will inevitably reduce the amount of tensile stress from other sources (service stresses, thermal stresses on

cooling, etc.) that the film is able to withstand before undergoing brittle failure (either cracking or delaminating) at some later epoch in the life of the system. These untoward stress conditions may be ameliorated by reducing the amount of film applied in any single coat [24]. Flat, platelike pigments also tend to reduce the amount of internal stress buildup in coating films [23]. These considerations become singularly more critical when the original adhesion of the coating to the metal is in any way compromised. This is one of the primary motivations for metal scarification and the establishment of a viable anchor pattern before the primer is applied. On smooth surfaces (particularly thin gauged substrates where flexing and thermal distortions are maximized), high film thicknesses should not be applied. In these cases (coil coatings, aircraft skins, container coatings, etc.) thin-film systems based on barrier primers (or inhibitive primers) are preferred to high-build barrier systems.

D. SACRIFICIAL COATINGS Perhaps the most effective corrosion control coatings are the zinc-rich primers [12, 35, 36]. As this type of composition requires intimate contact with the steel being protected, these coatings are used exclusively as primers and, although many types are self-recoatable, they can never be employed over other types of coatings. Protecting steel cathodically (Fig. 67.3), these primers require high volumes of zinc pigment in order to ensure an electrical pathway across the film (the anode of the artificial galvanic cell) as well as between the film and the substrate (the cathode). The zinc pigment employed in these coatings is in the form of a spherical dust of between 3 and 20 mm in diameter (usually 7 mm). Contact necessary to maintain current transfer is thus tangential (Fig. 67.4) between individual spheres of pigment and between the particles of pigment and the metal surface. In order to sustain the electrical continuity requirements of the primer, the spherical zinc dust particles must be packed sufficiently close together to ensure this contact [11]. This configuration allows for little binder, and formulations are quite critical in order to maintain electrical conductivity without depriving the film of the level of binder necessary to maintain adhesion and acceptable physical properties. Two fundamentally different classes of zinc-rich primer have been developed, the organic zinc-rich primer and the inorganic zinc-rich primer [36]. D1. Organic Zinc-Rich Primers Organic zinc-rich primers are based on a variety of resin systems, including epoxy/polyamides, high-molecularweight linear epoxies, moisture-cured urethanes, highstyrene resins, chlorinated rubbers, and epoxy esters, and the chemical and physical properties of the film in part reflect

SACRIFICIAL COATINGS

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FIGURE 67.3. Fundamentals of zinc-rich protection.

Conduction Across Thin Vehicle Sheath Between Zinc Dust Particles and Steel Substrate(Va + Vc)

Conduction Across Thin Vehicle Sheath Between Zinc Dust Particles at Tangential Contact Points (Va + Va)

Metallic Conduction Through Zinc Dust Particles

In Porous Films MONOMOLECULAR Conductivity may Occur VEHICLE SHEATH (Va) via Interstitial Electrolyte

ZINC DUST PARTICLE (VP)

ZINC PRIMER FILM

Interstitial Vehicle Steel Substrate

Metallic Conduction Monomolecular through Steel Vehicle Sheath (Vc) Substrate

FIGURE 67.4. Electrical conduction in zinc dust pigmented primer.

the resin system used. Such primers, employing an organic binder, are not, in fact, radically different from other organic paint films except for their very high volumetric concentration of pigment (65% zinc by volume of the total dry film volume). The formulation of organic zinc-rich primers has been treated at length by several authors [13, 37, 38]. In this type of primer, each individual sphere of zinc is lightly encapsulated with a monomolecular layer of binder that facilitates film cohesion and provides adhesion to the substrate. While it is necessary that all zinc particles bear at least a monomolecular layer of binder on their surface to ensure adhesion and film cohesion, it is important that the presence of the binder sheath around each particle of paint does not introduce levels of electrical resistance that are too high. Conductivity must be high enough to support the necessary unidirectional flow between cathode and anode. Control of

the volumetric ratio of conductive pigment to dry nonconductive binder ensures that the binder sheath around each particle of pigment (zinc) does not become so thick that the current transfer is too greatly reduced or so depleted that cohesion and film strength are lost [37]. Optimum loadings of zinc will actually depend on the specific formulation, the presence of auxiliary pigments (thixotropes, coloring agents, inhibitors, etc.), and the geometries of the pigment packing within the dried film [11, 13]. D2. Inorganic Zinc-Rich Primers The inorganic zinc-rich primers are fundamentally different from the organic [11]. Instead of employing a binder which encapsulates the zinc, these materials employ a reactive binder, usually an inorganic silicate (either an alkaline

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FIGURE 67.5. Stylistic representation of postulated structure of inorganic zinc-rich film in which silicate vehicle is primary valency bonded to zinc atoms on particles of zinc dust pigment. There is no encapsulation, and film is porous to ingress of electrolyte. This affords good film strength, adhesion, electrical conductivity, and cathodic protection.

silicate, such as sodium, potassium, lithium, or a quaternary ammonium silicate, or a partially hydrolyzed alkyl silicate, such as ethyl silicate) [35, 36]. Such silicate binders convert on mixing with zinc dust to hard, cohesive films that are virtually wholly inorganic in nature. Reactions involve loss of solvent or water and hydrolysis of the silicate to the reactive acid (silicic acid), which then undergoes self-condensation and chemically reacts with zinc atoms on the zinc dust surface, forming a primary valency bonded matrix of tetrahedrally linked (poly) silicon oxide bridges between adjacent particles of zinc (Fig. 67.5) [35]. Although the reaction paths are somewhat different depending upon the silicate type used (alkaline or alkyl silicate), the eventual structures of all films are probably quite similar. For specifics on the reactions involved, the reader is referred to [12, 35, 36]. (A bonding similar to that which occurs between the silicate and zinc is also thought to occur across the interface with the iron in the steel surface [35], although this has never been definitively proven.) The resultant film matrix is entirely inorganic, open, and much more porous than the organic zinc film. In practice, films may also contain unhydrolyzed material (alkyl silicate or alkaline silicate), which may remain within the film for months after the material has originally dried. Subsequent reaction of residual binder with carbonic acid from the atmosphere probably completes the conversion of residual silicate to silicic acid and adds to film cohesion. In service, the open film may become saturated with corrosive electrolyte, such as brine, which is readily absorbed into the film

porosities. This will also form zinc corrosion products that will further convert the film. The presence of brine solutions within the film will add conductivity to the film, and, as long as sufficient zinc remains to act as an anode, even in the presence of chloride solution, the primer will retain its ability to protect steel at least in the short term. Unless given the opportunity to dry out, however, zinc reaction products may be dissolved or eroded away leading to a more linear loss of zinc. Inorganic zinc-rich films are harder, stronger, and far more adherent than are the organic zinc-rich films [35]. As the film matrix is entirely inorganic, these primers also have better resistance to solvent and to heat than the organic zinc-rich primers and may be used for tank linings and applications involving temperatures as high as 400 C (750 F). The lack of encapsulating binder in inorganic zinc-rich primers means that zinc loading concentrations are not as critical to galvanic conductivity as they are in the case of organic zinc-rich primers, although reduced anode loadings will inevitably produce some reduction in performance. For the most part, performance levels realized from the inorganics are superior to those realized from the organics, although surface preparation and application requirements (especially application of the alkaline silicate-based materials) are more exacting. Also vulnerable is the security of intercoat adhesion between organic topcoats and the inorganic primers, especially in the case of the alkaline silicates. Alkaline silicates based on sodium silicate, in the absence of high baking temperatures, are cured by the postapplication of amine phosphate solutions, residues of which must be

INHIBITIVE PRIMERS

carefully removed if subsequent deadhesion of topcoats is to be avoided. D3. Postgalvanic Protection Phenomena In both organic and inorganic zinc-rich primers, the duration of cathodic activity is finite, and the effects of zinc polarization and the generation of zinc corrosion product gradually convert protection from the galvanic mechanism to one that is, at least in part, a barrier mechanism [11, 12]. The corrosion product blocks and seals the porosities of the film with a dense inorganic coat. In the case of the organic, this polarizing film is most usually a surface phenomena. In the more porous inorganic, zinc corrosion product occurs within the interparticulate spaces, building up and more completely sealing the film. Other postgalvanic phase mechanisms have also been suggested, including inhibition from zinc products [39], locally elevated pH, and control of oxygen reduction [40]. Most authorities, however, ascribe long-term protection to barrier effects [4, 11, 12, 38]. The interval of galvanic protection before conversion to the secondary protecting mechanism depends on the type and composition of system and the nature of the environment. Galvanic activity in alkyl silicates has been recorded after three years under atmospheric conditions. In epoxy zinc-rich primers, the cathodic protection phase is briefer [41].

E. INHIBITIVE PRIMERS E1. Pigmentation Inhibitive primers are typified by traditional compositions based on red lead and linseed oil, alkyd resins pigmented with zinc yellow (zinc potassium chromate), and, more recently, the many proprietary epoxy, alkyd, urethane, and latex systems that rely on modified phosphate, borate, and molybdate pigments [15, 42, 43]. While traditional chromate pigments are still being employed where inhibitive primers are used for the most demanding service (aircraft coatings, automotive primers, coil coating primers, etc.), the press of environmental and toxicological concerns threatens to eventually eliminate their employment in all coatings in the same manner that led to the removal of leadbased pigmentations. All of the above pigment types either have some direct limited solubility in water [44] or form reactive products (of limited solubility) with certain binders or their degradation products [45]. Several reviews of the many traditional and newer offerings of these types of pigments have been published [14, 15, 42, 46]. Nevertheless, the exact mechanism whereby these various inhibiting species inhibit corrosion remains imprecisely understood and is probably not the same for all inhibitors.

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It is, however, believed that chromates, and possibly molybdates, function by decreasing the oxidizing threshold at which passive films are naturally formed [44]. While oxygen alone will effectively passivate steel at high enough concentrations, those concentrations required for the passivation of steel (unlike the case with aluminum) exceed the levels of oxygen that dissolve in water at neutral pH. Oxidation levels high enough to induce passivation on steel are possible using oxidizing inhibitors, such as the chromates, however, which are thought to form passive films of complex chromic and ferric oxides on the metal [47]. Auxiliary pigmentations which increase the basicity of the film (zinc oxide, wollastonite, etc.) increase the pH of the interfacial environment, and at these levels of alkalinity, oxygen concentrations necessary for passivation fall nearer to those levels of oxygen that will dissolve in water [48]. It is thought, therefore, that these auxiliaries act in tandem with the inhibitor, reducing the amount of inhibitor necessary to achieve passivation under any given set of conditions. In many of the more popular binder systems for this type of product (e.g., alkyds, oils, and other fatty acid–based systems), these same basic pigments (both inhibitive and auxiliary) may react with acid groups on the binder, acting as pigmentary “crosslinking centers,” introducing unwanted embrittlement into the film. These effects must be controlled by judicious balancing of the formulation. Active ions from other inhibitors are thought to act to reinforce the naturally occurring oxide layer rather than establish a passive film themselves. Lead cations from lead-based inhibitors, as well as the modified zinc phosphate inhibitors, are thought to plug up discontinuities in the natural oxide layer, reinforcing it and thereby preventing ionic egress from the metal to the electrolyte. In spite of their long history, inhibition from lead-based pigments is itself incompletely understood. It is most generally thought to be attributed to the reaction of certain acidic moieties (azeleic acid, pelargonic acid, and other long-chained mono- and dicarboxylic acids) from oxidizing binders with lead monoxide giving soluble azelates and pelargonates which serve as a continuous source of inhibitive lead cations [45, 49]. Binders that produce higher concentrations of lower molecular weight acids (e.g., acetic acid), in lieu of the azelates, give much less effective protection [49], as do pigments such as lead dioxide [15]. These components appear to be corrosive rather than inhibitive. E2. PVC/CPVC and Controlled Permeability Inhibitive metal primers are designed with relatively high pigment volume concentrations. This design facilitates sufficient water absorption into the film, so that soluble inhibitive ions may be released by the pigment (or its salts with the binder) and carried to the metal surface beneath

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FIGURE 67.6. Corrosion protection by inhibitive primer.

the primer where passive films may be established. Formulation devices are designed so that enough water is able to penetrate the film and dissolve sufficient inhibitor to provide ionic concentrations high enough to achieve passivation at the subfilm metal (Fig. 67.6). The permeability to ionic solutions must not, however, be so great as to allow depassivating ions (chloride, sulfates, etc.) into the film from the environment. These depassivators raise the inhibitor threshold necessary to achieve passivation [50]. Where depassivator concentrations are high enough, it may not be possible to achieve passivation with even the strongest inhibitor. Fortunately, in normal primer films, the relative ionic impermeability of the coating films and the presence of the source of passivating ion within the film encourage favorable passivator/depassivator ratios at the substrate and preserve protection. Selection of the pigment/binder composition should be such that the inhibitor is not so soluble or the film so permeable that excessive release of soluble inhibitive ions leads to the premature exhaustion of the film’s available inhibitive reservoir and, in consequence, loss of primer effectiveness [51]. (Additional dangers with the use of very soluble inhibitors in paint films are the propensity of such films to exhibit osmotic blistering in high-humidity and freshwater conditions, especially when inhibitive primers are employed beneath lower permeability finishes.) Alternatively, the binder should not be so impermeable (or the pigment not so insoluble) that insufficient inhibitive ions are released to establish or sustain protection. This condition will lead to little effectiveness as an inhibitive primer,

although such systems may make rather inefficient barrier films. There remain unanswered questions concerning these mechanisms, for it is difficult to understand how passivating films may form on metals bearing tightly adherent organic films. It has been suggested that inhibition may not become established until some initial adhesive breakdown in a quasibarrier protection has occurred [32]. E3. Water-Based Inhibitive Primers Recent emphasis on minimizing volatile organics released from the application of coating films has encouraged the use of water-based systems. With some important exceptions [9], the films deposited using such technologies are more water sensitive than are those deposited from solvent-based systems. Thus the films do not generally make good barrier coatings. They do, however, make more suitable binders for inhibitive primers. Water-based primers of this type are based on acrylic latexes, epoxies, and water-borne alkyds. In these systems, the inhibitor exists in contact with the water carrier in the can prior to the application of the paint, and this constitutes a special case. Here, the inhibitive ions are already in solution at the time of application, and, where passivating films may form on the steel before the wet primer film dries, the initial passive film may be established without the access of water through the coating in the field [51]. Thus, classical inhibition from the access of external water to the pigment in service is presumably necessary only to sustain the passive film.

SURFACE PREPARATION

F. SURFACE PREPARATION Although the primary foci of this chapter are the mechanisms by which coating films may be employed to control metallic corrosion, such discussions cannot be considered complete without some reference to the preparation of the metal surface prior to coating application. The relevance of good surface preparation to the effectiveness of applied coating systems on steel has been well recognized since the work of Hudson in the 1930s, 1940s, and 1950s [52]. Since then, numerous other researchers have reaffirmed the singular importance of good practice in this regard, whether we are considering the sandblasting of structural steel, the anodizing of aluminum, or the pickling of nickel surfaces [53–55]. Morcillo [56] emphasizes the particular importance of good surface preparation to the performance of modem corrosionresistant coating systems on steel. The contribution of surface preparation to coating system performance may be reduced to one or more of the factors described in the following paragraphs. F1. Removal of Interference Material The presence of numerous and diverse conditions and contaminations on practical metal surfaces is virtually axiomatic. Oils, greases, and other organics may reduce the surface energy of normally high-energy metal surfaces from values near 400 dyn/cm to values 65 C (l50 F) can cause most adhesives, which some coatings use as the main bonding mechanism, to flow and move toward the lower parts of the structure. When this happens on pipe, a loose shell of the outer material is left on the top half of the pipe with no bond to the pipe. Many tapes, shrink sleeves, and extruded polyolefins are susceptible to this problem. This process may take only a few months to several years to occur, according to the temperature and the type of coating. If water penetrates this space, corrosion becomes a problem. When high temperatures are encountered, polypropylenes instead of polyethylene are normally used. Coal tar and asphalt based coatings can slowly flow with gravity, leaving the top part of the structure exposed. Fusion-bonded epoxy (FBE) coatings can absorb more water than normal at higher temperatures but do not flow or move on the structure. If undertfilm contaminants are present, FBE coatings can blister or disbond, as shown in Figure 68.3. B5. Surface Preparation It is widely recognized that surface preparation is the most important single factor in coating performance [2]. Many in the coating industry agree that two-thirds the cost of any coatings job should be spent on surface preparation of the structure to be coated. One should always prepare the best surface possible for a given coating situation. Blasting using

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FIGURE 68.3. Blistering of FBE coating on hot pipeline (180 F). (Photo furnished by Lone Star Corrosion Services, Lancaster, TX.)

the correct material and method not only cleans the metal surface but also provides a surface profile (anchor pattern) for coating adhesion. Other methods of surface preparation include the use of hand tools, power tools, water blasting, and wire brushes. The Steel Structures Painting Council and NACE International are both excellent resources for information on surface preparation in the coating industry. Field repairs or coating replacements do not always allow for the best surface preparation. When the proper surface preparation is not possible, replacement coatings are selected that are more tolerant to poor surface preparation. Coating tests performed on a variety of prepared surfaces, as well as past experiences with coatings that performed well in a particular condition, should be used to help select the best coating for the situation. For field repairs, leaving the structure uncoated (if adequate Cathodic Protection (CP) is available) may be a better choice than to apply a coating that may not bond to the structure because of poor surface conditions. This is especially true during cold or damp weather or if the structure is sweating or cold. Hot applied coatings should never be applied on a cold structure. Petrolatum-based coatings are sometimes used on cold or wet structures. If a structure is left uncoated or if a poor coating condition exists, this part of the structure should be monitored and properly coated when the conditions are more suitable. Most plant-applied coatings require more stringent surface preparations. Surface preparations are much easier to control and perform in plant setting, as shown in Figure 68.4. Some of the common surface preparation mistakes made in plants are: Using the wrong type of blast material Using contaminated blast material Improperly using surface treatments Using rinse water that contaminates the steel surface

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SELECTION AND USE OF COATINGS FOR UNDERGROUND OR SUBMERSION SERVICE

FIGURE 68.4. Inspection for and grinding of surface defects after blasting in FBE. (Photo furnished by EB Pipe Coating, Inc., Panama City, FL.)

FIGURE 68.5. Proper handling of FBE coated pipe with padded forks. (Photo furnished by EB Pipe Coating, Inc., Panama City, FL.)

Allowing foreign materials, such as dust and debris, to accumulate on the structure surface Using air that may be contaminated with oil or water; debris, tape, and other contamination on conveyors, wheels, and supports

Field applications are by hand or machine but are usually labor intensive. Applying coatings in the field requires training and practice with the particular coating and application method chosen. Tape coatings for pipelines appear to be easily applied. Too many times, tape coatings are applied over the wrong or uncured primer with improper tension or not enough overlap. Shrink sleeves can be over- or under heated, be unevenly heated, leave air pocket, or allow debris to be blown between the sleeve and the pipe. Brushable or sprayed coatings must be allowed to properly cure before applying the second coat or before the structure is back-filled or submersed. Frequently, these coatings are applied too thin or thick. Liquid coatings must be mixed properly, especially two-part coatings, such as epoxies. Urethanes and polyurethanes are very susceptible to moisture and humidity and require specialized equipment and training for application. Any coating applied in the field must be compatible with any existing coating. The transition from one coating to another is a very critical area. Many coating failures have occurred at the transition area.

Some coating processes require that a chromate and water mixture be applied to the structure before the corrosion coating is applied. Chromate is a conversion coating that reacts with steel to change the surface chemistry. Too much chromate can cause a buildup that may affect the coating bond. Weather conditions should always be considered whether in the field or a plant. Flash rusting after surface preparation but before coating application can occur if the humidity is too high. Flash rusting can cause some coatings to fail. Oil or grease contamination must be removed with proper solvents. Blasting will only spread the oil or grease on the structure surface. The original surface condition of the structure must be considered. New pipe or structural steel may have mill lacquers, mill scale, or other surface contaminants. Stored or in-service structures may have considerable corrosion, scale, or other contamination. Leaking product can contaminate surfaces of in-service pipes or tanks in leak locations. B6. Application Methods Field and plant application techniques vary considerably for each coating system. Plant applications are usually much faster, and each step can be controlled more easily. The coating itself is usually applied with automatic equipment on a production line. Inspections and more sophisticated testing can easily be performed. Storage, handling, and transportation can present a problem for plant-coated structures, as shown in Figure 68.5.

B7. Weld Area, Joint, and Additional Component Coatings Whether the main coating is plant or field applied, there are areas around girth welds (called field joints on pipelines), flanges, bolts, and other joined areas where the structure is connected together that must be considered when choosing the primary coating. In some cases, a very good coating is used on the main structure, but then an inferior or poorly applied coating is used in the areas where connections are made in the field. Field connections are the most critical area on many systems. Heat-affected zones and other stressed

CRITERIA FOR SELECTION OF COATINGS FOR UNDERGROUND OR SUBMERSED STRUCTURES

areas around welds can corrode quickly when exposed to an electrolyte. Bolted or riveted connections can have crevices and other areas that can easily corrode or cause stress points where cracking can occur. The best field connection coating for a particular system is usually the same as or better than the coating used on the main structure. Any field connection coating must be compatible with the previously applied coating. The field connecting process should cause very little or no damage to the existing coating. Prequalification testing of the applicator will help to ensure the crew is capable of applying the coating as specified. B8. Repairs Any coating selected must be easily repaired with a material that is compatible with, or the same as, the parent coating [3]. Many times repair coatings are improperly applied, not allowed to cure, or incorrectly mixed. Many of the problems mentioned above for field connections and field-applied coatings are the same for repair coatings as well. Once again, testing and experience are important when choosing a repair coating. B9. Cathodic Protection Since corrosion always requires the presence of an electrolyte (moisture) in contact with the metal, if a metal could be coated with a material which was absolutely waterproof and absolutely free from holes, all attack would be stopped [4]. Always assume “there is no perfect coating”! Even if it were possible to apply and install a “perfect” coating, deterioration, soil stress, environmental factors, and damage from outside forces would soon cause portions of the coating to fail. For this reason, cathodic protection must be applied and properly maintained. This synergistic relationship is well proven and documented. Cathodic protection can affect all coatings. The hydrogen, hydroxyl ions, and other electrochemical reactions caused by cathodic protection currents at the cathode may cause blistering or cathodic disbondment of the coating. Hydroxyl ions are one of the most aggressive chemical species and nearly all organic binders are capable of reacting with them [5]. Cathodic protection can cause loss of adhesion between the coating and the metal, leaving a void between the coating and the structure. The surrounding environment and the level of cathodic protection at a particular site determine the extent of cathodic disbondment. The amount of cathodic protection needed to cause cathodic disbondment is very difficult to determine. There are several rules of thumb but little data to actually use for determining the level that affects a particular coating in a certain environment. The amount of cathodic protection needed to achieve adequate protection varies greatly for each type of coating. The effective electrical strength of a coating is often

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expressed as the resistance per average square foot of coating [6]. Determining the actual amperes per square foot needed to protect the coated pipe after installation can provide valuable information for future projects. There are charts, tables, and computer programs used for determining the amperes of current per square foot needed for a particular coating system. Current requirements or current density measurements are related to coating conductance [7]. One must also consider the coating breakdown factor. How long will it take the coating to deteriorate before cathodic protection has to be increased? At 10, 20, or 30 years after installation, what percent of the structure will be exposed because of deteriorated or damaged coating? Coatings that incorporate an electrically insulating outer layer, such as polyolefin tape coatings, can preclude effective CP from reaching the pipe surface in disbonded areas [8]. Some research and considerable debate have not fully answered the question of whether cathodic protection can be provided under disbonded coating. There is some information available that indicates some protection can be achieved when the electrolyte present under the disbonded coating has very low resistance (e.g., saltwater environments). The use of pulsed cathodic protection has also been studied to determine its benefits (if any) for providing cathodic protection under disbonded coatings. When selecting a coating system to be used with cathodic protection, the “nonshielding” or “fail-safe” characteristics of that coating may be more important than other issues that are normally considered [8a]. Nonshielding means if the coating system adhesion fails and water penetrates, corrosion on the metal is significantly reduced or eliminated when adequate CP is available. Fusion-bonded epoxy is a nonshielding coating. B10. Handling, Storage, and Transportation of Coated Pipe Coating materials selected must be able to withstand the rigors of handling, storage, and transportation for a particular project. Coated pipes and other structures should be handled with padded forks, hooks, or other equipment that will not damage the coating. Stacking and storage should be according to structure weights, with the proper separation materials between each piece. The separation material should not damage the coating in any way. When coated structures are stored outside, separation helps to protect the coating by not letting debris settle in openings and crevices. During transportation, the coated structures must be properly strapped down with soft straps, not chains or bands that may damage the coating. Coated structures should also be loaded and fastened in a manner that does not allow the coated structures to shift or slide during transportation. Once again, proper separation and support should be used between the structures. Rocks and other road debris can damage

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coatings during transportation. Mud flaps and protective shields help eliminate most of these problems during truck or rail transportation. Outside storage can cause problems with all coatings. Ultraviolet (UV) rays damage most materials with time. Polyolefin coatings become brittle and crack. The FBE coatings can chalk. This chalking is usually only 1 or 2 mils in depth, but rain can wash off the “chalked” layer and expose new FBE to the sun. Coal tar coatings tend to crack with exposure. Heat and cold can cause some coatings to contract or expand or cause the adhesive to flow. B11. Specialty Coating and Overcoats Coatings may be specific for a particular situation. One may involve weight coatings. Concrete is applied over coated pipe to provide enough weight to sink the pipe or provide it with mechanical protection. Other “overcoats” are applied to provide mechanical protection when pipe or other structures are pulled through road or river bores, put inside casings, or installed in rocky terrain or for other unusual circumstances. Other special “overcoat” material may be used in splash zone areas of subsea structures. These overcoats can shield the cathodic protection if they disbond from the metal surface. If the reinforcement in concrete coated pipe or structures contacts the pipe, the cathodic protection will protect the reinforcement but not the pipe.

C. COATING TYPES FOR UNDERGROUND OR SUBMERGED STRUCTURES This discussion covers the variety of coatings available to industry at this time. The information provided about each system comes from numerous tests, specification writing and reviews, coating inspections, field applications and evaluations, literature reviews, and plant applications. Each coating being considered for use on an underground or submerged system should be evaluated using the above criteria and any other criterion that may affect the performance of the coating. C1. Coal Tar Based Coatings C1.1. Application. Coal tar based coatings were one of the first types to be successfully used as a corrosion coating. These coatings are made of coal tar mixed with various blends or formulations of fillers and extenders. As with most types of coatings, there are several grades of coal tar coatings available. One must study the different types available to ensure the best choice. Coal tar mixtures have been widely used on steel underground storage tanks. Most are applied at the tank manufacturing facility by spraying or brushing the coal tar epoxy or mastic onto clean, primed surface. Surface

preparation normally involves brushing or blasting. A reinforcing wrap may be used to add mechanical strength. These tanks are not always carefully handled during transportation and construction. Coating damage is easily repaired by cleaning the area and applying a coal tar based epoxy or mastic. Coal tar coatings have also been successfully used on subsea structures, such as offshore platforms and pipelines, submersed bridge structures, docks, and other submersed metal structures. Plant-applied coal tar enamels have been used for many years on underground pipelines, Coal tar is applied to a blasted and primed pipe surface. The surface is usually grit blasted in accordance with SSPC-SP6 after any dirt, oil, grease, and so on, are removed. The pipe is preheated to remove any moisture and keep the steel temperature above the dewpoint. A heated [50 C (120 F) or less] type A or type B primer is applied by spraying as the pipe rotates down the conveyor line. The type A primer is made from coal tar. Type B is a fast-drying synthetic and is normally specified. The coal tar normally used for plant application contains fully plasticized enamels. Hot coal tar enamel flows onto a prepared pipe as it rotates down a conveyor. After the coal tar flows onto the pipe, a felt wrap is applied to reinforce the coating. This is normally topped with whitewash or kraft paper to provide ultraviolet protection. The thickness of the 3 coal tar and 1 felt wrap is normally 2:4 mm 32 in: ð0:8 mm 32 in: Þ Early felt wraps contained asbestos, but in the 1980s most of the industry changed to felt that is fiberglass reinforced. There are other methods for plant application of coal tar, but this method is the most common. Coal tar coatings are not as popular as they once were but are still a major part of the coating industry. Variations of coal tar enamels, epoxies, mastics, and urethanes are used for field applications and repairs. Coal tar enamels are hot-applied materials that are usually applied with wraps containing coal tar coated fiberglass felt. Coal tar epoxies are solvent-cured coatings that are usually brushed or hand applied in multiple layers. Coal tar urethanes are multicomponent applied mixtures that are normally used for large-scale rehabilitation projects. Coal tar mastics are hand applied and are used for a variety of irregular shaped structures. C1.2. Pluses. Coal tar enamel coatings have a long history of corrosion protection. Many coal tar coated pipelines have been in service for over 50 years and are still in very good condition. Coal tar coatings are easily repaired with fieldapplied coal tar enamel, coal tar epoxy, or tape coatings. Coal tars are thick coatings that have excellent electrical insulating properties, have low water permeation properties, resist bacterial attack, and have the solvent action of petroleum oils. C1.3. Minuses. Even though many pipelines are well coated and still perform as intended, many pipelines have major

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coal tar coating failures. Much of the failed coal tar was applied over the ditch, on pipe surfaces that were not well prepared, or in undesirable weather conditions. In many cases, plant-applied coal tar was not applied using an acceptable method. Handling during transportation, storage, and construction also contributed to these failures. If the temperature is not properly maintained, both the application and future performance of coal tar enamel coating are compromised. Coal tar epoxy is a multicoat application that requires time for solvents to evaporate before burying or submersing. Coal tar urethanes require sophisticated equipment and well-trained individuals for application. Application of urethanes is not recommended in highmoisture situations. Coal tar enamel coatings are subject to soil stress that may cause the coating to wrinkle, crack, disbond, and expose steel surfaces. If water penetrates under the disbonded coating, cathodic protection shielding can become a problem. Cathodic protection requirements normally increase as the coal tar coatings age. Operating temperatures of pipelines coated with coal tar are normally limited to 65 C (150 F) or less, Coal tar enamels can “cold flow,” leaving the top of the pipe without adequate coating, especially on high-temperature pipelines. There are safety concerns with hot-applied coal tar enamels. Not only are there the problems associated with the dangers of hot application, but, in addition, the fumes may be toxic. The early use of asbestos felt has been a problem for those removing and disposing of these coatings. Each of these problems can be studied and handled with proper education and precautions. C1.4. Improvements. During the 1990s, there have been improvements in coal tar enamel coatings, including an effort to make these coatings more tolerant to higher temperature operations. C2. Fusion-Bonded Epoxy C2.1. Application. Fusion-bonded epoxy pipeline coatings have been used successfully since the late 1960s. A typical FBE formulation consists of epoxy resins, curing agents (hardeners), catalysts and accelerators, prime and reinforcing pigments, control agents (for flow and stability), and specialty ingredients [9]. The first FBEs were applied over a primer, but later developments allowed them to be used without a primer. Originally, FBEs were applied much thinner than the 12–16 mils normally specified in 1999. The FBE has been used to coat rebar used in bridge, road, and building construction to help prevent corrosion of the rebar in concrete. The application process for FBEs is one of the most stringent and complicated in the industry. Because of the thin film and potential for water absorption, the surface

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cleanliness must be very good. The pipe is moved through the plant on a series of inspection racks and conveyors. Oil, grease, or other contaminants are removed from the pipe. The pipe is then preheated to keep it above the dewpoint during the surface preparation phase. After blasting to SSPC-SP10 (Near White), most companies now require a phosphoric acid wash and rinse to further clean, etch, and provide surface energy to the pipe that helps to attract and bond the FBE to the steel surface. Many companies (especially in Europe) require the application of a chromate solution for an additional surface treatment. After the surface preparation, the pipe is heated to a range of 232 C (450 F) to 260 C (500 F) by either heat induction coils or gas-fired heaters. The FBE powder is applied to the hot pipe by a dry air spray system. Electrostatic gun are used to attract the optimum amount of powder to the pipe surface. The FBE powder melts, gels, and cures as the pipe rotates down the conveyor line. Gel times are usually in a range of 10–30 s at these application temperatures. Once the powder has gelled, the FBE on the pipe is hard enough to support the weight of the pipe on the conveyor tires. After the FBE is cured (usually 30–100 s) a water-quenching system cools the coated pipe. The coated pipe is then inspected for acceptance, rejected, or placed on hold for necessary repairs. C2.2. P1uses. The FBE coating has the best bond (adhesion) to steel of any pipe coating. The bond is mechanical as well as chemical. The FBE is very flexible and allows for field bending of pipelines or rebar. The FBE mostly cures by the time it is in the quench, so it can be handled immediately after the process. Repairs are made easily with two part epoxies or patch sticks. The water absorption of FBEs is a plus when disbondment occurs. Unlike thick coatings that have high electrical resistance, FBE (because of the water absorption) will allow enough cathodic protection current through the film to protect the pipe under the coating. Therefore, corrosion and pitting are rarely encountered under disbonded FBEs if adequate cathodic protection is provided. Unless severe failure occurs on an FBE coated structure, the cathodic protection requirements normally do not increase significantly as the FBE ages. Stress corrosion cracking has been studied extensively and has never been observed on FBEcoated pipelines [10]. The FBE coatings handle well during transportation and construction activities. Comments have been made by some in the industry that they do not use FBE because it is easily damaged during construction, compared to thick coatings. When FBE coatings are damaged, the damage is easily found and repaired. The thicker coatings may “hide” damage that has occurred and disbondment may not be seen in some cases until the pipe fails. Multilayer coatings with FBE as the base or primer coat have been very successful when properly specified and

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applied. These coatings have been used successfully on pipelines with internal operating temperatures up to 150 C (300 F) or pipelines in critical areas. These systems are discussed further in Section C3. C2.3. Minuses. The FBEs require more stringent surface preparation and application techniques than most plantapplied coatings, thus requiring more attention during the coating process. Slivers and other steel imperfections can cause “holidays” that may not be a problem on thicker coatings. Field joints are expensive if FBE is used for the field joint coating (even though FBE is the preferred coating). Higher temperatures can cause water to be absorbed more quickly and cause disbondment if surface contaminants are present. For pipelines operating at temperatures between 65 C (150 F) and 85 C (l85 F), thicker, up to 800 mm (32 mils), FBE should be used. Flexibility during bending can be a problem when using thicker FBE coatings. When stored outside, UV rays can cause the FBE to “chalk” and become grainy in the first few mils of coating. Multilayer coatings with FBE as a base coat are expensive but offer an excellent coating. C2.4. Improvements. Recent changes in FBE coatings include the development of FBEs that will perform better in high-temperature operations. Each manufacturer offers several formulations of FBE. Each formulation has specific uses and they usually vary in gel, cure time, application technique, and test results. The FBE-based multilayer systems have become an important part of the FBE market. New formulations of FBEs are being tested that can be applied and cured at much lower temperatures to reduce the amount of energy required and to accommodate the higher strength steels that may be affected by higher application temperatures. C3. Polyolefin Coatings Plant-applied polyolefin coatings have been used for pipeline coatings since the early 1960s. Because of application techniques, plant-applied polyolefins are not normally used on other structures. Since polyolefins do not bond well to steel, it is normally applied over an adhesive or other product that provides the bond to the steel. Polyethylene has been the polyolefin of choice for most extruded coatings, but polypropylene is normally used for higher temperature [>65 C (150 F)] operations. For many extruded polyolefins, an adhesive that is usually a rubber-modified asphalt adhesive or a butyl rubber compound provides the bond to the steel. Three-layer systems use an epoxy primer or a coating quality FBE for a base coating to the steel; then an ethylene copolymer or terpolymer adhesive is used for the “tie” layer. Some of these multilayer systems are now using chemically modified polyolefins for the topcoat or the tie layer. These

chemically modified polyolefins chemically and mechanically bond to the FBE base coat. C3.1. Polyolefin over Adhesive. Polyolefins essentially have no adhesion to steel [11]. When adhesives are used as the bond to steel, the process usually involves heating the pipe to a temperature above the dewpoint. The pipe is then blasted and the heated adhesive is applied by flood coating onto the pipe. One method uses a crosshead extrusion die system that extrudes the melted polyolefin over a rubbermodified asphalt adhesive coated pipe as it travels down a nonrotating conveyor system. Another method uses a rotating conveyor system and side extruders to apply the adhesive and the polyolefin. After preheating and blasting, the pipe rotates down a conveyor where the molten butyl rubber adhesive is spirally applied by the side extruder. The molten polyolefin is then side extruded over the adhesive. This process allows the layers of adhesive and polyolefin to fuse together and produce a theoretically seamless pipe coating. C3.2. Polyolefin over Epoxy Primer. Three-layer systems using a liquid epoxy primer or powder epoxy primers also use a rotating conveyor system for the application process. These primers are sprayed onto the hot pipe as it rotates down the conveyor. The co-terpolymer is then side extruded over the primed pipe as the “tie” layer between the epoxy and polyolefin. The polyolefin is then side extruded over the co-terpolymer. This system has been used in Europe and the Middle East for over 20 years with excellent results. The three-layer coating evolved from the extruded two-layer (co-terpolymer adhesive plus polyethylene) polyethylene coating first used in Europe in about 1960 [12]. The addition of the epoxy primer improved the cathodic disbondment characteristics of these coatings. Many threelayer systems now use powdered epoxies or FBE for the primer coating. C3.3. Polyolefin over FBE. Today, the trend for many multilayer coatings is to use a thicker 300–625 mm (12–25-mil) FBE base coat. The FBE is then topped with a chemically modified polyolefin (CMP) as either the topcoat or the middle layer of a three-layer system. This is an improvement over the traditional three-layer systems because of the thicker FBE quality base coat and use of the CMP. The bond of the CMP to the FBE is normally much stronger than the bond of the co-terpolymer to the FBE. The bond of the CMP to the FBE is both chemical and mechanical and will not allow a separation between the two when properly applied. When a three-layer system is used, the hot CMP will bond well to the top layer of polyolefin, making a coating system that is excellent. Some of these coating systems now use a powdered CMP for the middle layer of a three-layer system


or top layer of a two-layer system. These have been very successful coatings when applied properly. C3.4. Pluses. Plant-applied polyolefin coatings are overall excellent coatings. Polyolefins are among the best materials to prevent water permeation. The “slick” surface of plantapplied polyolefins is not normally affected by soil stress. These systems handle easily during transportation and construction activities. Polyolefins are resistant to many chemicals, environments, and bacterial attack. The polyolefin over FBE systems is an excellent pipeline coating system for hightemperature service. Test results show that coating systems consisting of polypropylene (PP) over FBE can be used successfully when internal operating temperatures reach 150 C (300 F) [13]. Even though polyolefin over FBE coating systems can be more expensive than other coatings, they provide an excellent overall coating system when specified and applied properly. C3.5. Minuses. One of the problems with polyolefins over adhesive coatings is that the adhesive layer provides the only bond to the steel. Adhesive can allow the polyolefin to shrink and expand on the pipe surface as the temperature changes. If this movement occurs, steel can be exposed or it can cause problems with field joint coatings. If the polyolefin is damaged while in service, it can disbond or split and allow water penetration between the polyolefin and pipe (Fig. 68.6.) Since the polyolefin has very high electrical resistance, cathodic protection is shielded. Underfilm corrosion can then become a significant problem, since the primer does not offer much corrosion protection. External damage can allow water to permeate the polyolefin more easily even though the film is not broken enough for the holiday detector voltage to penetrate. This external damage can also break the bond between

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the adhesive and the polyolefin. This is especially true of the crosshead die extrusion process. The co-terpolymer tie layer for some polyolefin over epoxy primer coatings also allows the top layer to shrink and expand as the temperature changes. Many times, the epoxy primer used is very thin, 80–130 mm (3–5 mils), and may not provide corrosion protection if the top layers are damaged. If water were to penetrate between the polyolefin and the primer, the thick polyolefin would shield the cathodic protection. One of the major disadvantages of all polyolefin coatings is coating selection for use on repairs, field joints, valves, and other components. Most tapes and shrink sleeves have problems with soil stress, do not bond well to the polyolefin, or have application problems that may cause failures at the repair or field joint. It is difficult to coat bends and other components with most tapes and shrink sleeves. Liquid epoxies, mastics, and other such coatings do not bond well to the polyolefin, possibly allowing water to penetrate at the junction of the two coatings. One system coats the bends and many of the other components with an FBE covered with a powdered chemically modified polyolefin. This system works well with the polyolefin over FBE two-and three-layer systems. Since high temperature is required for this application, the heat can affect the plant-applied coating if this process is used for the field joints. These coatings are expensive when compared to most other plant-applied coatings. C3.6. Improvements. Recent developments include providing better repair and field joint coating systems for the three-layer systems. The use of thicker, high-quality FBE for the first layer and chemically modified polyolefins as the tie layer has improved the overall quality of these coatings. Flame spraying of thermoplastics and other materials is being developed to provide another method of applying these coatings to field joints, bends, and other components. C4. Tape Coatings

FIGURE 68.6. Damage to extruded polyethylene coat with adhesive only. Also note the movement of the coating from the field joint area. (Photo furnished by Lone Star Corrosion Services, Lancaster, TX.)

Several types of hand- and plant-applied tapes are used for corrosion protection on underground or submersed metals. Though some systems were developed earlier, most were developed in the late 1950s and early 1960s. Most have been used on buried piping systems, but some have been used for other applications. Tapes are normally applied over a primed surface, but some companies are promoting a “primerless” tape that does not require a primer. Polyethylen-backed tape systems have been the most widely used type of tapes in the pipeline industry. Other types of tapes are made from coal tar with special fibers, petrolatum, polyvinyl chloride (PVC), or polyolefin fiber mesh with compound. The proper selection and use of tape product depend greatly on the environment (e.g., soil conditions), size of pipe, and operating temperature.

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Polyethylene systems were developed from electrical and general use industrial tape applications. These tapes normally consist of one layer of polyethylene backing and an adhesive layer bonded to the polyethylene backing. The polyethylene backing layer in most cases is a low-density or blend of high- and low-density polyethylene. Adhesive compounds usually consist of elastomer butyl rubber, natural rubber, rubberized bitumen, or coal tar derivatives. Processing oils, fillers, tackifiers, and stabilizers are added to provide adhesion, shear resistance, and thermal and chemical resistance. Pigments and stabilizers are added to the polyethylene to provide color, UV resistance, and thermal stability and to improve aging. Polyethylene tapes come in several thicknesses of polyethylene and adhesive compounds. These tapes are usually made by one of three different methods: extrusion, coextrusion, or calendering. A polyethylene film is formed in a continuous sheet and the adhesive is laminated to it as the system is produced in large rolls. These rolls are then cut into the required widths and lengths. Some tapes use a woven polyolefin geotextile fabric for the backing. This product has several advantages over the solid polyethylene backed tapes that will be discussed later. These tapes use the same type of compounds for adhesion to the metal. The woven fabric provides the mechanical protection and the compound provides the corrosion protection in these coating types. Another type of tape is composed of coal tar base coating material supported on a fabric of organic or inorganic fibers. The fabric is covered on both sides by the coating materials. These tapes must be pliable enough to unwind from the roll during application. They are applied to the structure by heating the tape (usually with an open flame) on the structure side surface until the tape becomes liquid. The tape is then wrapped or applied to the primed surface. Many tape systems require that an outer wrap be placed over the tape product for buried service. These outer wraps may help to alleviate soil stresses on the coating and provide additional mechanical protection or insulation. In all these cases, outer wraps should be used that do not shield cathodic protection from the metal if the coating is damaged. C4.1. Pluses. Tape coatings are used in the field for repair, replacement, or base coatings. Most tape products can be quickly back-filled or immersed after application. Ease of application is another advantage of most tape products. Tapes can be applied by hand or with a tape machine. Even though there are a variety of tape types and prices, tape products usually cost less per square foot to apply than other coating products. Plant-applied systems are fast and handle easily. When properly selected for the environment and properly applied, tape coatings can provide an economical, easily applied coating system.

Tape systems can be applied in several layers to provide a thick coating when needed for certain environments. Tape coatings are normally very flexible and can be formed to fit many irregular shapes. Tape coatings normally have high electrical resistance (or low conductance) and therefore low cathodic protection requirements, especially for the polyethylene backed tapes. C4.2. Minuses. Tape systems have been widely used for coating pipelines and other structures, but there have been major failures with the solid film backed tape products. Soil stress, poor surface preparation, and poor application techniques have been reasons for most tape coating failures. Other major reasons for tape failures have been the use of an incorrect primer, not using primer when required, or not allowing the primer to cure properly before tape application. There have been cases of bacterial attack on some butyl rubber compounds used for the adhesive. At this time, tapes have a very limited use on structures that operate at temperatures >65 C (l50 F). Soil stress effects on geotextile backed tape is much less than on most polyethylene backed tapes. The geotextile fabric has approximately a 20% stretch compared to the polyethylene backings, which can have up to a 600% stretch. When soil compresses around the pipe and forces downward on each side of a pipe, the force can cause the tape to wrinkle. Electrical resistance of polyethylene coatings and their susceptibility for unbonded installation create a serious problem on pipelines [14]. The electrical shielding comes primarily from the polyethylene backing. Even with the compound present, if water were to penetrate between the polyethylene backing and the metal, corrosion can occur because cathodic protection currents are shielded from the metal. On the woven fabric backed tape, soil can compress the compound and possibly move it to expose the metal surface, but the woven fabric will not shield cathodic protection currents. C4.3. lmprovements. Recent improvements with tape coatings involve the use of materials that will function effectively at higher temperatures. The success and technology of the fiber mesh backed tapes provide a superior product because their stretch is minimal, and therefore, soil stress is not as significant a problem. The woven mesh backing will not shield cathodic protection currents even if the compound is compressed and the metal is exposed. One company has developed a mesh backed tape that is nonshielding to cathodic protection should there be a disbondment and water penetrate between the coating and the pipe. After 20 years of use this coating system has had no reports of corrosion under the coating as with other types of coatings, such as solid film backed tape and shrink sleeves, even when improperly applied [14a].


C5. Shrink Sleeves Shrink sleeve coatings are normally made from cross-linked heat-shrinkable thermoplastic backing (usually polyethylene) that serves as a tough permanent outer layer. Because of radiation cross linking these materials have an elastic memory that allows the product to be supplied in an expanded state. When heated, the material shrinks. Similar technology is used in the electrical connector industry. Special adhesives, from soft sealants to highly crystalline hot melts, are applied (similar to tapes) to the backing material. This technology was developed in the early 1960s and continues today. These products were introduced to compete with coldapplied tapes. Early shrink sleeve applications had serious adhesion problems. Later versions have major improvements in adhesion and application procedure. Even with recent improvements, application procedures for shrink sleeves are very critical to their performance. During inspection, poor adhesion at the interface dual epoxy/sleeve has been detected mainly because of the insignificant amount of heat applied to the sleeve [15]. There are basically two types of shrink sleeve on the market. One is a sleeve that is installed over a pipe end and slid into place. This type requires the sleeve to be installed before making connections. The other type is a wrap-around sleeve with a closure strip that is normally preattached to the sleeve for easy application. The closure has an adhesive to help hold the sleeve in place until it is shrunk to the structure. These sleeves have the advantage that they can be applied on any structure (normally pipes) even after construction. C5.1. Pluses. Shrink sleeves are quick and relatively easy to use for field joint and repair coatings. They provide a tough, durable coating when properly selected and applied. Shrink sleeves are compatible with a variety of plant-applied coating systems. They work best when used on new construction or systems where the metal surface can be heated to the desirable application temperature before the shrink sleeve is applied. C5.2. Minuses. As mentioned above, shrink sleeves must be properly applied. Even though the process appears to be easy, instructions must be followed. The most frequent mistakes involve not heating the metal surface properly before application or not using the proper heat to apply the sleeve. Metal on in-service pipelines that are operating at low temperatures or structures in cold weather regions may not allow for proper metal heating before applying the sleeve. If the structure is too cold, the adhesive next to the metal may not melt enough to allow proper adhesion, even though the sleeve will shrink and appear to be properly installed. Improper heating usually occurs because the wrong type of torch is used or certain areas are heated either too much or not enough. Overheating can cause the sleeve to overshrink and

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possibly cause the sleeve to crack or split. If not heated sufficiently, the sleeve may not properly bond to the metal. Shrink sleeves can have the same problems with soil stress, cathodic shielding, and temperatures >65 C (150 F) as polyolefin backed tapes. C5.3. Improvements. Shrink sleeves are being developed for use at high temperatures, which will make them easier to apply and less affected by soil stress. The use of infrared heaters for applying shrink sleeves has improved the process by allowing for more uniform heating of the sleeve and pipe. C6. Wax Based Coatings Waxes have been successfully used for coating materials on a variety of structures. Petrolatum tape systems are made from a combination of saturated petrolatum impregnated with a neutral compound covering a synthetic fabric carrier for strength. Another tape uses pure wax in combination with synthetic fibers. As with most tapes, the proper primer must be used. Wax based tapes are easily formed to fit irregular shapes. Hot applied wax is normally composed of microcrystalline wax, fillers, and sometimes wetting agents and corrosion inhibitors. These coatings are applied by first melting the wax at temperatures in a range of 94–260 C (200–500 F) and flood coating the structure. An outer wrap is normally applied over the wax to provide mechanical protection to the pliable wax material. Another coating is a hand-applied coal tar wax mastic that typically consists of microcrystalline or petrolatum compounds. After hand application, the mastic is usually covered with a protective wrap to help prevent mechanical damage. C6.1. Pluses. Wax coatings are easy to apply and work well on irregular surfaces. When selected and applied properly, these coatings can provide excellent corrosion protection. C6.2. Minuses. Wax based coatings are soft and sensitive to high temperatures. As with any hot-applied coating, there are safety concerns. C7. Asphalt Asphalt coatings were derived from petroleum and have been used similarly to coal tar coatings on a variety of structures. Asphalt coatings were cheaper than coal tar coatings, but they did not have the same chemical resistance and bendability and were not effective at higher temperatures. Application was similar to that of coal tar. Asphalt mastics are used on a variety of irregular structures. They are normally hand applied, and, if allowed to cure properly, provide a good coating system for certain environments. Asphalt coatings are not widely used at this time for underground or submersed metal corrosion protection.

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C8. Liquid Coatings Liquid coatings for underground or submerged structures are very diverse but must be selected with caution. Thus a coating has worked well above ground or in certain types of submersion service does not mean this same coating will perform well in conjunction with cathodic protection when buried or submersed. Extensive testing should be performed to ensure that the coating will provide the required protection in that environment. These coatings can be applied in plants or at construction sites to a variety of structures, valves, and other components that are not easily coated by other coating methods. Many of these coatings are applied in the field for replacement or repair coatings. The application methods include, but are not limited to, brushing, air spray, airless spray, and plural component spraying. Since the list is extensive, only the most popular and most promising will be discussed. C8.1. Epoxies. Coal tar epoxies, as mentioned above, have been used for pipeline, tank, and structure coatings for many years. They have been fairly successful coating but are not flexible, easily crack, and are affected by temperatures over 65 C (150 F). There are several types of 100% solids epoxy coatings formulated for burial or submersed use in conjunction with cathodic protection. These are excellent coatings when properly applied and used in the correct environment. These coatings are fast-cure, high-build materials with good adhesion. Epoxies of 100% solids are environmentally safe and do not shield cathodic protection to the extent that thicker coatings do. Some of these epoxies are formulated to perform well on systems operating at temperatures up to 100 C (212 F). Epoxies must be mixed in the correct ratio for the coating to cure properly. C9. Urethanes and Urethane Blends Urethane and urethane blends are becoming more popular as coatings for burial or submersed structures. These coatings have been successfully used over FBE coatings to provide mechanical protection during construction, in rocky areas, and for road and river crossings. The bond to the metal is not as strong as some other liquid coatings, but improvements are being made. These are very fast curing, 100% solids that are normally applied with plural component spray equipment. Urethanes used for underground or submersion service are normally adversely affected by moisture or high humidity. C10. Epoxy Phenolic One coating that has been used in other industries for years is modified epoxy phenolic [16]. Epoxy phenolics have been successfully used in some applications, such as the internals of

tanks and other systems that operate at temperatures up to 121 C (250 F) with some success. Cathodic protection may damage some of these coating applications at these temperatures. More study and field experience are required at this time. C11. Polyurea Polyurea coatings are new to the underground and submersion coating industry. At this time, most of these are modified polyurea systems. They are very rapid cure even at temperatures less than 29 C ( 20 F) but have high thermal stability, even at temperatures greater than 150 C (300 F). Polyureas are moisture insensitive and have high abrasion resistance. These systems are normally applied using plural component, high-temperature/high-pressure impingement mix application equipment. At this time, polyureas have not performed well in cathodic disbondment and moisture permeation testing, but these coatings do hold promise if the adhesion to the metal surface can be improved. C12. Esters Esters are normally used as polyester or vinyl esters. They are used in areas where high temperature is a concern. Esters must be applied to a very well prepared surface with a very good anchor pattern to have adequate adhesion and good cathodic disbondment and moisture permeation results. This material is normally applied in two coats with relatively short recoat time intervals. Mixing must be monitored closely since large component mixing ratios (up to 64 : 1) are normally used.

D. TESTS TO EVALUATE UNDERGROUND OR SUBMERSION COATINGS There are many different standard tests for evaluating and comparing the various coatings mentioned above. The tests most often used and some of the various organizations that provide the standards are given in this section. Frequently, these standard tests are modified to provide test information for particular environments or conditions. Short-term laboratory testing is limited in providing longterm performance information but is very valuable in providing information about differences between various types of coatings. The longer the test, the more meaningful the information. Shortterm tests are performed at coating plants during production to determine if the application process is providing the specified product. Sample preparation for each type of test is very important. When possible, samples should be taken from the actual production or field application. When this is not possible, the application process should match the production or field application as much as possible. One should not rely totally on one set of test results to select a coating system, because all

SHIELDING VERSUS NONSHIELDING COATINGS

997

CSA (Canadian Standards Association) SSPC (Steel Structures Paint Council) API (American Petroleum Institute)

FIGURE 68.7. Cathodic disbondment test of a tape coating system. This particular system failed the test (Photo furnished by ITI Anti-Corrosion, Inc., Houston, TX.)

the performance parameters of a coating system cannot be obtained using only one type of test. One must consider all the various conditions of application and service to select the test criteria for a particular coating system. Tests normally performed for all underground or submersion coatings: Cathodic disbondment test (CDT) as shown in Figure 68.7 Moisture permeation (sometimes called hot water soak or adhesion testing) Resistance to impact Flexibility (sometime called bend test) Chemical resistance Adhesion (strength of bond of the coating to the metal) Fusion bonded epoxies are additionally tested for: Porosity (through film and interface) Surface contamination (amount of contamination on the back side of the coating surface) Thermal characteristics by differential scanning calorimetry (DSC) Tape coatings are additionally tested for: Adhesion or peel strength Shear, stretch, or soil stress characteristics One should be familiar with the various organizations that publish standard procedures for testing coatings. Some of these are: NACE International ASTM (American Society for Testing and Materials)

The most valuable information is that taken from inspections performed on in-service coatings. Any time a coating is exposed after being in service for a length of time, the coating system should be inspected and tested using many of the same tests mentioned above. One should inspect for blisters, disbondment, adhesion strength, discoloration, and any other adverse effects to the coating. Results of cathodic disbandment tests are presented in Table 68.1. When a coating is used in conjunction with cathodic protection, the pH should be checked under any blisters or disbondments. This will provide information on the effectiveness of the CP system to provide protection in these areas. This is discussed in more depth in the next section.

E. SHIELDING VERSUS NONSHIELDING COATINGS When using coating with cathodic protection, it is critical for the end user to consider what potential problems could exist if the coating system fails and electrolyte penetrates between the coating and the metal being protected. Shielding of the cathodic protection current by disbonded external coatings on pipelines causes more external corrosion than any other process in today’s pipeline industry. The resulting corrosion cells cannot be protected by the cathodic protection system, so corrosion will develop, as illustrated in Figure 68.8. There have been many articles written about the problems of external corrosion caused from disbonded and shielding pipeline coatings [17–23]. When selecting a pipeline coating, the nonshielding characteristics of the coating system may be more important than other issues normally considered. To adequately protect underground pipelines, a coating must conduct CP current when disbondment occurs [24]. Nonshielding means if the coating system adhesion fails and water penetrates corrosion on the metal is significantly reduced or eliminated when adequate CP is available, Fusion-bonded epoxy is a nonshielding coating (Fig. 68.9). Corrosion and stress corrosion cracking (SCC) have not been an issue under disbonded FBE in over 40 years of service or the mesh backed tape after over 20 years of service (Fig. 68.10). Other coatings may also provide this property but should be studied and tested to ensure this property. This has been proven by the use of the various in–line inspection (ILI) tools that are used to find corrosion and other defects in pipelines. With the use of the external corrosion direct assement (ECDA) methods now being used, companies can evaluate

998


TABLE 68.1. Cathodic Disbondment Test Results Comparing Tape Coatings and Shrink Sleevesa System

30 days

60 days

90 days

1. Two-layer tape/primer 2. Shrink sleeve/two-part primer 3. Shrink sleeve/no primer 4. Shrink sleeve/two-part primer 5. Tape/primer 6. Tape/primer 7. Shrink sleeve/two-part primer 8. Tape/two-part primer 9. Shrink sleeve/two-part primer 10. Shrink sleeve/primer 11. Shrink sleeve/primer 12. Mesh-backed tape 13. Two-layer tape/primer 14. Two-layer tape/primer 15. Two-layer tape/primer

41 mm 25 mm 0 mm >50 mm Complete >50 mm Complete 8.3 mm Complete Complete Complete 0 mm 10.7 mm 38.3 mm >75 mm

45 mm >65 mm 5.6 mm >50 mm Complete >50 mm Complete 6.3 mm Complete Complete Complete 0 mm 12.3 mm >50 mm >75 mm

58 mm >65 mm 0 mm >50 mm Complete >50 mm Complete 8.9 mm Complete Complete Complete 0 mm 12.5 mm >50 mm >75 mm

failure failure failure failure failure



a

Provided by Lone Star Corrosion Services, Ennis, TX.

FIGURE 68.8. Shrink sleeve applied in 1997 and resulting corrosion found in 2006. Shrink sleeve shielded the CP. Pipe potential met all NACE criteria.

the condition of the external coatings. When doing the ECDA diagnosis of any failed coating and potential corrosion, one of the most critical evaluation methods is to check the pH of any water under the coating. The electrochemical process of cathodic protection will cause an alkaline area where the current enters the metal. External corrosion may occur for other reasons, such as AC or DC interference or shielding from other materials that may shield the CP. The cause of external corrosion must be properly evaluated to distinguish between shielding and other causes.

FIGURE 68.9. Water under blisters on FBE coated pipe used for gas transmission in central United States. Water under the blisters had a pH of 12 with no metal loss showing nonshielding properties of FBE. (Photo furnished by Lone Star Corrosion Services, Ennis, TX.)

If the coating system is nonshielding to the CP current, the pH will typically be 9 or above, indicating that CP is being effective even under the disbonded coating. At this pH level, corrosion on steel surfaces will be limited or nonexistent since steel is protected at these levels of pH. The high pH also indicates adequate CP is available under the coating. Disbonded coatings that are shielding CP current will typically have a pH of 7 or less. Many times significant corrosion will develop under these coatings because CP cannot be effective under these coating systems, yet the CP potentials taken above ground will meet or exceed the various CP criteria.

REFERENCES

FIGURE 68.10. Mesh backed tape applied to pipe with condensation on one side. After three years in service, adhesion was very good on the dry side, but there was no adhesion on wet side, but the pH of the water under the coating was 11 with no corrosion on the pipe, proving the nonshielding properties of this coating system. (Photo provided by Polyguard Products, Inc., Ennis, TX.)

F. CONCLUSIONS There may be other selection criteria that are important for a particular situation, but the ones given in this chapter are always important. Testing of materials and judgment of experienced individuals will help to determine the correct coating for a particular structure and situation. Clear, precise specifications are very important. Specifications should be continually updated to reflect changes in the industry. It is essential to have well-trained and qualified inspectors who have a real passion for ensuring the best coating possible under the circumstances. If coatings are plant applied, preproduction meetings, coating, and testing help to ensure that the plant understands the specifications and can produce coated material as specified. For field-applied coatings, well-trained applicators should be employed. Once again, preproduction meetings, coating, and testing are advised. Company employees who apply coatings should be trained for each particular coating type to be used. Many factors must be considered when selecting a coating for a particular system. Short-term testing will help to determine which coatings will perform best in certain conditions. Past experience with a particular coating system used in the same environment should always be considered. Good coating performance in one environment does not mean the same coating will perform well in a different environment. Once selected, the coating must be properly applied to a clean, well-prepared surface. The coated product must then be handled, transported, and stored properly during

999

the construction phase. After installation, the parameters of the coating, such as temperature and cathodic protection limits, should not be exceeded. Well-written coating specifications and inspection in the field or plant by inspectors who are well trained and passionate about what they are doing will help to ensure the best possible coating system under the circumstances. When selecting a pipeline coating, the shielding and nonshielding characteristics of the coating system may be more important than other issues normally considered. Most of the external corrosion on pipelines today is caused from coating that have disbanded (for a variety of reasons), allowed water to penetrate between the coating and the pipe, and then shield the CP current allowing corrosions to develop. More CP does not solve this problem; only recoating the pipe properly will solve the problem.

REFERENCES 1. A. C. Coates, “Pipeline Recoating—A Cover Up Story,” Pipeline Digest, Apr. 1991, p. 11. 2. Good Painting Practice, Steel Structure Painting Manual, Vol. 1, 3rd ed., Steel Structure Painting Council, 1993, p. 19, 3. R. Norsworthy, “Select Effective Pipeline Coating,” Hart’s Pipeline Digest 34(4), 17 (Feb. 1997). 4. Corrosion Basics, An Introduction, National Association of Corrosion Engineers, NACE International, Houston, TX, 1984, p. 213. 5. G. B. Byrnes, Mater. Perform., Sep. 1989. 6. Control of Pipeline Corrosion, National Association of Corrosion Engineers, (NACE International), Houston, TX, 1967, p. 18. 7. J. L. Banach, “Evaluation Design and Cost of Pipe Line Coatings,” Pipe Line Ind., Mar. 1998, p. 62. 8. D. A. Diakow, G. J. Van Bovan, and M. J. Wilmott, Mater. Perform., 37(5), 17, (1998). 8a. R. Norsworthy, “Is Your Pipeline Coating ‘Fail Safe,” Pipeline Gas J., 233(10) 62 (Oct. 2006). 9. T. A. Pfaff, “FBE Serve A Broad Market,” Hart’s Pipeline Digest, Oct. 1996, p. 20. 10. D. Neal, “Fusion-Bonded Epoxy Coating: Aging and BelowGround Performance,” Hart’s Pipeline Digest, 34(13), 20 (Sep. 1997). 11. D. Neal, Pipeline Gas Ind., 81(3), 43 (1998). 12. S. J. Lukezich, J. R. Hancock, and B. C. Yen,“State-of-theArt for the Use of Anti-Corrosion Coatings on Buried Pipelines in the Natural Gas Industry,” Topical Report, Document No. 92/0004, Gas Research Institute, Chicago, Apr. 1992. 13. R. Norsworthy, “High Temperature Pipeline Coating Using Polypropylene Over Fusion Bonded Epoxy,” in International Pipeline Conference, Vol. 1, ASME, New York, 1996.

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14. C. D. Tracy, Pipeline Gas Ind., 80(2), 27 (1997). 14a. C. Hughes and R. Norsworthy, “Proven Protection,” World Pipelines, Oct. 2007, p. 49. 15. V. Rodriquez, E. Perozo, and E. Alvarez, Mater. Perform., 37(2), 44 (1998). 16. R. Norsworthy, Pipeline Gas J., 225(3), 44 (1998). 17. J. A. Beavers and N. G. Thomson, “Corrosion Beneath Disbonded Pipeline Coatings,” Mater. Perform. Apr. 1997, p. 19. 18. “Coatings Used in Conjunction with Cathodic Protection,” NACE International, Technical Committee Report, Item No. 24207, Houston, TX, 2000, p. 2. 19. S. Papavinasam, M. Attard, and R. W. Revie, “External Polymeric Pipeline Coating Failure Modes,” Mater. Perform., Oct. 2006, p. 28.

20. F. M. Song, D. W. Kirk, D. E. Cormack and D. Wong, “Barrier Properties of Two Field Applied Pipeline Coatings,” Mater. Perform., Apr. 2005, p. 26. 21. J. A. Beaversand N. G. Thompson, “Corrosion Beneath Disbonded Pipeline Coatings,” Mater. Perform., Apr. 1997, p. 13. 22. G. R. Ruschau and Y. Chen,“Determining the CP Shielding Behavior of Pipeline Coatings in the Laboratory,” Corrosion 2006, Paper No. 06043, NACE International, Houston, TX, 2006. 23. J. Alan Kehr, Fusion Bonded Epoxy (FBE)—A Foundation for Pipeline Corrosion Protection, NACE International, Houston, TX, 2003. 24. D. P. Moore, “Cathodic Shielding Can Be a Major Problem After a Coating Fails,” Mater. Perform., 39(4), 44 (2000).

69 ENGINEERING OF CATHODIC PROTECTION SYSTEMS J. H. FITZGERALD III Grosse Pointe Park, Michigan

A. B. C. D.

Introduction Operation of cathodic protection Uses of cathodic protection Types of cathodic protection D1. Galvanic anodes D2. Impressed current D2.1. Power sources D2.2. Anode E. Galvanic anode design E1. Typical uses E2. Advantages and limitations E3. Design process E3.1. Design parameter E3.2. Current requirement E3.3. Anode selection E3.4. Anode requirement E3.4.1. Cathodic protection circuit resistance E3.4.2. Anode output E3.4.3. Number of anodes required E3.5. Anode life E4. Grouped anodes E5. Anode installation F. Impressed current design F1. Typical uses F2. Advantages and limitations F3. Design process F3.1. Design parameters F3.2. Current requirement F3.3. Anode selection F3.4. Anode requirements and life F3.5. Groundbed resistance F3.5.1. Anode-to-ground resistance F3.5.2. Cable resistance F3.5.3. Structure-to-ground resistance F3.5.4. Total circuit resistance F3.6. Rectifier selection F4. Other groundbed configurations

G. Design of other cathodic protection systems G1. Marine structures G2. Heat exchangers G3. Steel in concrete G4. Water storage tanks G5. Above-ground storage tank bottoms H. Other design concerns H1. Electrical continuity H2. Shielding H3. Cathodic interference H4. Attenuation I. Criteria for cathodic protection I1. Steel and cast iron I2. Aluminum I3. Copper References

A. INTRODUCTION This chapter presents the practical application of cathodic protection. Various uses are shown along with the principles of design for galvanic anode and impressed current systems. This chapter also presents recommended steps and formulas as well as overall information that will guide the engineer in the choice and design of the most appropriate cathodic protection for the structures involved. Standards and other publications containing information on various facilities are referenced. Engineering of cathodic protection involves not only the calculation of current requirements, resistances, and voltages but also an understanding of the type and configuration of cathodic protection to be used. The designer must consider its


1001

1002

ENGINEERING OF CATHODIC PROTECTION SYSTEMS

practicability, life-cycle costs, and maintenance and operational requirements. B. OPERATION OF CATHODIC PROTECTION Essentially, cathodic protection involves the application of a direct current (dc) from an anode through the electrolyte to the surface to be protected. This is often thought of as overcoming the corrosion currents that exist on the structure. That is not really what happens as there is no flow of electrical current (electrons) through the electrolyte. There is, of course, a flow of ionic current in the electrolyte as explained in earlier chapters on electrochemistry. Cathodic protection eliminates the potential differences between the anodes and cathodes on the corroding surface. A potential difference is then created between the cathodic protection anode and the structure such that the cathodic protection anode is of a more negative potential than any point on the structure surface. Thus, the structure becomes the cathode of a new corrosion cell. The cathodic protection anode is allowed to corrode; the structure, being the cathode, does not corrode.

D1. Galvanic Anodes Galvanic anode protection is often called “sacrificial” because the anode is thought of as “sacrificing” itself to protect the structure. This type of protection utilizes a galvanic cell consisting of an anode made from a more active metal than the structure. The anode is attached to the structure, either directly or, to permit measurement of the anode output current, through a test station. Magnesium and zinc are the most common galvanic anodes for underground use. In salt water, zinc anodes and aluminum alloy anodes are commonly used. In freshwater, magnesium is frequently used. For underground use, magnesium anodes are packaged in a backfill consisting of 75% gypsum, 20% bentonite, and 5% sodium sulfate. The purpose ofthe backfill is to absorb products of corrosion and to absorb water from the soil to keep the anodes active. Magnesium and zinc are also available in ribbons and extruded rods. Galvanic anodes require no external power. The protective current comes from the electrochemical cell created by the connection of the anode material to the more noble or electrically positive metal of the structure. D2. Impressed Current

C. USES OF CATHODIC PROTECTION Cathodic protection is widely used to protect many structures. Among these are underground structures, on-grade tank bottoms, marine facilities and ship hulls, and water storage tanks. It can also be used for stray-current corrosion control. See Kumar [1] and [2] and [3, Chapter 6–8] for further information. Effective cathodic protection may or may not require electrical isolation of the structure. Where the protected structure is not electrically isolated, protective current requirements are increased as some current will be lost to interconnected facilities. Many structures such as gas distribution piping, transmission lines, and small piping systems can usually be easily isolated. Other facilities such as large tank farms, wharves, foundation or sheet piling, and complex piping networks often cannot be isolated economically or with any degree of certainty. The designer must consider the feasibility of isolation in the overall current requirements and in the layout of cathodic protection. The presence or absence of isolation may also affect the type of cathodic protection that the designer chooses. D. TYPES OF CATHODIC PROTECTION There are two types of cathodic protection: galvanic anode and impressed current. Typical uses and selection recommendations are discussed later.

Impressed current protection provides dc from a power source. The current is delivered to anodes made of a material having a very low or essentially inert dissolution rate. The anodes serve simply to introduce the protective current into the electrolyte. D2.1. Power Sources. The most common power source for impressed current protection is the transformer rectifier. This unit, commonly called simply a rectifier, reduces incoming alternating current (ac) voltage and rectifies it to dc. There are also solid-state “switchmode” rectifiers that perform similar functions without the use of transformers. Rectifiers can be provided with constant voltage, constant current, or structure-to-electrolyte potential control. In areas where electrical power is not readily available, solar power– and wind-driven generators coupled with storage batteries are used, There is also some use of thermoelectric cells, in-line turbine generators (in gas or oil pipelines), and internal combustion engine–driven generators. D2.2. Anodes. A variety of materials are used for impressed current anodes, Among the oldest are high silicon, chromium-bearing cast iron, graphite, and junk steel. Magnetite and lead–silver anodes are also used, with lead–silver being confined to use in seawater. Among newer materials are “dimensionally stable anodes,” so-called because the anode itself consists of a deposit on an inert substrate. This deposit may be consumed, but the

GALVANIC ANODE DESIGN

anode shape tends to remain stable. Included in this category are platinized niobium or titanium and mixed-metal oxide (MMO)/titanium anodes. Underground impressed current anodes are usually backfilled in a carbonaceous material such as metallurgical or calcined petroleum coke. The purpose of the backfill is to increase the effective size of the anode, thus reducing its resistance to earth, and also to provide a uniform environment around the anode, increasing its life. Anode life is extended by the large coke backfill column since the current is discharged from the coke column as opposed to being discharged from only the anode. Another advantage to increasing the size of the anode is that the resultant reduction in anode current density reduces acidity in the vicinity of the anode.

E. GALVANIC ANODE DESIGN E1. Typical Uses Galvanic anode systems are typically used where protective current requirements are relatively low, usually in the range of several hundred milliamperes to perhaps 4 or 5 A. Offshore structures, having a current requirement of many hundreds of amperes can also be protected by large galvanic anodes weighing anywhere from 135 to 635 kg. Common places to use galvanic anode protection include well-coated, electrically isolated structures, offshore structures, ship hulls, hot-spot pipeline protection, heat exchanger water boxes, and in environments of resistivity below 10,000 W cm, although they can be used in much higher resistivities, especially if coating and isolation conditions are favorable. E2. Advantages and Limitations There are advantages and limitations to galvanic and impressed current systems. The designer needs to assess the engineering and economic aspects of each making the type of protection system to use: Advantages of galvanic systems . No external power required . Little maintenance . Relatively easy installation . Little chance of cathodic interference . Less inspection and recordkeeping than with impressed current systems Limitations of galvanic systems . .

Lack of adjustment without resistors in anode circuits Limited current output

. . .

1003

Possible high replacement costs Need for good coating Need for electrical isolation of protected structure

E3. Design Process The following will serve as a guide to the designer for the steps to follow during the design process. The example is based on an underground pipeline. Later in this chapter the reader is referred to publications and standards that will provide information on other structures. Refer also to [5], Chapter 6] for additional information on galvanic anode design. E3.1. Design Parameters. To begin the design, establish the electrolyte resistivity, the protective current requirement, the desired life of the anodes, whether or not the structure is electrically isolated, whether or not there are any straycurrent concerns, and the physical configuration desired for the anode. E3.2. Current Requirement. There are two ways of establishing current requirement. If the structure is in place, one can test it as shown in Figure 69.1. In this procedure, a test current is applied to the structure and the resultant change in structure-to-electrolyte potential is measured. From the data, the current requirement can be calculated. Details on current requirement testing are given by Peabody [4] and in [5, Chapter 5]. In new construction, the current requirement is often calculated from estimates such as those in Table 69.1. The amount of bare steel depends on the assumed quality of the coating, both to begin with and after several years of operation. A coating efficiency, for example, of 95%, equals 5% bare metal. The current requirement is then the total area of the pipe times the percent bare (as a decimal) times the current per square meter, or Ireq ¼ A %bare mA=m2

ð69:1Þ

Where Ireq is the total, current requirement in milliamperes and A is the total area of the structure in square meters. E3.3. Anode Selection. Table 69.2 lists the characteristics of several types of galvanic anodes. There are two types of magnesium anode alloys, standard (H-1) alloy and highpotential alloy. Generally speaking, high-potential anodes are desirable if the electrolyte resistivity exceeds 8000 W_cm. There are two grades of zinc anodes, one for seawater use and one for underground use. There are also two grades of aluminum alloy anodes for saltwater use. If zinc anodes are chosen, it is important that they meet the purity required by the American Society for Testing

1004


FIGURE 69.1. Typical current requirement test setup. (From Advanced Course Text, Appalachian Underground Corrosion Short Course [5]. Used by permission.)

TABLE 69.1. Typical Current Requirement for Cathodic Protection of Bare Steela Environment

(mA/m2)

Neutral soil Well-aerated neutral soil Highly acid soil Soil supporting sulfate-reducing bacteria Heated soil Stationary freshwater Moving, oxygenated freshwater Seawaterb

4.5–16.0 21.5–32.0 32.0–160.0 65.0–450.0 32.0–270.0 11.0–65.0 54.0–160.0 32.0–110.0

a Data from Air Force Manual 88–9, Corrosion Control, Chapter 4, p. 203, and [1, Table A-1]. b May be as high as 160.0–430.0 in cold and Arctic waters.

and Materials (ASTM) standard B-418 Type I (MILL-Spec A-18001) for seewater use and ASTM B-418 Type II for underground use, Lower purity anodes will not function properly. Zinc anodes are generally limited to environments of resistivity below 1500 W cm because of their low driving voltage. Zinc anodes are used successfully in higher resistivity soils, however, on some very well coated, electrically isolated facilities such as underground storage tanks. E3.4. Anode Requirement. This step involves calculation of the resistance of the cathodic protection circuit, the potential difference between the anode and the structure and, from these numbers, the anode output.

TABLE 69.2. Galvanic Anode Characteristica Consumption Material

Theoretical Output (A-h/kg)

Zinc type I type II

860 816

Actual Output (A-h/kg)

Efficiency

Rate (kg/A-year)

Potential to CDEb

781 739

90% 90%

11 12%

1.06 1.10

Magnesium H-1 alloy

2205

551–1279

25–58%

6.8–16

1.40–1.60

Magnesium High Potential Al/Zn/Hg Al/Zn/In

2205 2977 2977

992–1191 2822 2591

95% 87%

7.3–8.6 3.1 3.3

1.70–1.80 1.06 1.11

a Data adapted from Advanced Course Text. Chapter 6. Appalachian Underground Corrosion Short Course, West Virginia University, Morgantown, WV, 2008, and from [3, p. 139]. b Copper–copper sulfate electrode.

1005

GALVANIC ANODE DESIGN

E3.4.1. Cathodic Protection Circuit Resistance. The resistance of a vertical anode to ground can be calculated from the following equation (after H. B. Dwight [6]), based on the dimensions of the anode package (Eq. (69.2) [4, p. 134] and on accompanying computer disks): r 8L ln 1 ð69:2Þ Rv ¼ 2pL d where Rv is the resistance to earth in ohms, r is the soil resistivity in ohm centimeters, L is the anode length in centimeters, and d is the anode diameter in centimeters. If the designer decides to use horizontal anodes, the resistance to earth is different from the vertical anode and is given by (see [6] and [1, pp. 1–8]) qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi3 0 2 2 4L þ 4L ð2hÞ2 þ L2 r @ 4 5 ln Rh ¼ 2pL 2dh

þ

2h d

qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi ð2hÞ2 þ L2 L

1

ð69:3Þ

1A

TABLE 69.3.

Copper Cable Dataa

Size AWG

dc Resistance 20 C W/1000 m

Max. dc Current Capacity (A)

14 12 10 8 6 4 3 2 1 1/0 2/0 3/0 4/0 250 MCM

8.4650 5.3152 3.3466 2.0998 1.3222 0.8334 0.6595 0.5217 0.4134 0.3281 0.2608 0.2070 0.1641 0.1388

15 20 30 45 65 85 100 115 130 150 175 200 230 255

a Data adapted from Advanced Course Text, Chapter 5, Appalachian Underground Corrosion Short Course, West Virginia University, Morgan town, WV 2008.

potential of the Structure (Vgs vs), or DVg ¼ Vga Vgs

where Rh is resistance to earth in ohms, r is the soil resistivity in ohm centimeters, L is anode length in centimeters, d is anode diameter in centimeters, and h is depth below the surface to the center of the anode in centimeters. The designer will find anode dimensions given in catalogs prepared by anode manufacturers. The packaged dimensions are used in the calculation. The anode is conected to an underground structure with a lead wire. The resistance of the wire (Rw) is very small but should be considered. Wire and cable data appear in Table 69.3. The resistance of the structure (Rs) to ground mayor may not be significant. In many cases, this resistance is small enough to be omitted from the calculation. Peabody [4, p. 135] provides additional information on calculating structure to ground resistance. The total cathodic protection circuit resistance (Rt in ohms) then is Rt ¼ R a þ R w þ R s

ð69:4Þ

where Ra is the resistance of the anode to ground from Eqs. (69.2), (69.3) or (69.9), Rw is the wire resistance, and Rs is the structure-to-ground resistance. E3.4.2. Anode Output. The anode output depends on the anode circuit resistance and the potential difference between the anode and the structure (the driving potential). The potential difference between the anode and the structure, DVg in volts, is the difference between the open-circuit potential of the anode (Vga vs,) and the desired polarized

ð69:5Þ

The current output, I in milliamperes, is then calculated from Ohm’s law: I¼

DVg 1000 Rt

ð69:6Þ

E3.4.3. Number of Anodes Required. This is calculated from the current requirement divided by the individual anode output: Anodes required ¼

current requirement anode output

ð69:7Þ

Anodes are usually spaced evenly along or around the structure. It is also good practice to place anodes near isolating fittings, building walls, or other locations where an inadvertent, although likely high-resistance, contact to another structure might occur. E3.5. Anode Life. Calculate the life of the anode from [5], L¼

Th W E UF hI

ð69:8Þ

where L is life in years, Th is the theoretical Ah/kg output (Table 69.2), W is the anode weight in kilograms, E is current efficiency (Table 69.2), UF is the utilization factor, h is hours per year (8766), and I is anode output in amperes. The utilization factor is usually chosen as 0.85 (85%); this means that once the anode is 85% consumed, its resistance to earth

1006


FIGURE 69.2. Typical galvanic anode system installation. (Courtesy of Advanced Course Text, Appalachian Underground Corrosion Short Course [5]. Used by permission.

begins to increase to the point that its output is reduced significantly. If the calculated life is insufficient, then the designer needs to choose a heavier anode or perhaps use more anodes than the design calculations require. Also, Eq. (69.8) can be solved for total required anode weight for a given life. The total weight can then be divided the individual anode weight to obtain the required number of anodes.

r 8L 2L ln 1þ ln 0:656N Rn ¼ 2pNL d S

ð69:9Þ

where Rn is resistance of the anodes to earth in ohms, r is soil resistivity in ohm centimeters, L is anode length in centimeters, d is anode diameter in centimeters, N is number of anodes, and S is center-to-center spacing of anodes in centimeters.

E4. Grouped Anodes There are times when the designer may desire to group several anodes together in a bank. The anodes may then be connected to the structure through a test station. This is particularly useful when replacing anodes to minimize the number of excavations required. When placed in groups, the individual anode outputs are reduced due to a mutual interference resistance proportional to the anode spacing. This resistance decreases as anode spacing increase. The resistance to earth of a group of vertical anodes may be calculated from the Sunde equation (see [4, p.134] and [7]):

E5. Anode Installation Figure 69.2 shows a typical installation. It is good practice to use test stations wherever possible. Test stations are necessary to evaluate the effectiveness of cathodic protection on underground structures; placing at least some anodes at test stations permits measurement of the output current and “instant off ” polarized potentials at that location. Knowledge of anode current aids in evaluating anode performance and in calculating eventual anode life. Figure 69.3 shows anodes on a distribution pipeline.

FIGURE 69.3. Typical galvanic anode cathodic protection.

IMPRESSED CURRENT DESIGN

F. IMPRESSED CURRENT DESIGN

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structure is electrically isolated, and the physical configuration desired for the groundbed.

F1. Typical Uses Impressed current cathodic protection has wide application. It is especially applicable when current requirements are large, in some cases as high as 500 A or more. Then, too, a small impressed current system putting out less than an ampere might be used to replace a dissipated galvanic anode system. Common uses of impressed current include long transmission pipelines, complex underground structures, pilings, marine structures and ship hulls, replacement for dissipated galvanic systems, large condenser water boxes, reinforcing steel in concrete, bare or poorly coated structures, unisolated structures, and water storage tank interiors. F2. Advantages and Limitations As discussed in Section E, there are advantages and limitations to galvanic and impressed current systems. The designer needs to assess the engineering and economic aspects of each in making the choice of the type of protection system to use: Advantages of impressed current systems . Adjustable output . Large current available . Applicable to poorly coated or bare structures . Applicable to nonisolated structures . Low-cost method of replacing spent galvanic anodes Limitations of impressed current systems . .

. .

Constant power required More maintenance, inspections, and recordkeeping than with galvanic systems, especially for regulated tanks or pipelines Experienced electrical personnel may be needed for installation Possibility of cathodic interference

F3. Design Process The following will serve as a guide to the designer for the steps to follow during the design process. This example is based on an underground pipeline. Later in this chapter the reader is referred to publications and standards that will provide specific information on other structures. Refer also to [5, Chapter 5] for additional information on impressed current design. F3.1. Design Parameters. To begin the design, establish the electrolyte resistivity, the protective current required, the desired life of the groundbed (anodes), whether or not the

F3.2. Current Requirement. The current requirement can be determined by testing if the structure is in place as explained under galvanic anode protection and as shown in Figure 69.1 and further described by Peabody [4]. In new construction, the current requirement is often calculated from estimates such as those in Table 69.1. The amount of bare steel depends on the assumed quality of the coating both to begin with and after several years of operation. A coating efficiency, for example, of 95% equals 5% bare metal. As described in Section E3.2 for the current requirement in a galvanic anode system, the current requirement for an impressed current system is also calculated by multiplying the total area of the pipe, the percent bare (as a decimal), and the current per square meter, or Ireg ¼ A % bare mA= m2

ð69:1Þ

where Ireq is the total current requirement in milliamperes and A is the total area of the structure in square meters. F3.3. Anode Selection. The designer has a variety of anodes from which to choose and with experience will learn which anodes perform best for the situation at hand. High silicon, chromium-bearing cast iron, and graphite have similar characteristics. Dissipation rate varies with the environment, but 0.5 kg/A-year is typical; steel dissipates at 9.1 kg/ A-year. Kumar et al. [1] present data on MMO anodes. For characteristics of other anode materials, refer to the manufacturer’s literature. F3.4. Anode Requirements and Life. In determine the number of anodes required for an impressed current groundbed, it is best to start with the calculation of anode life. While the dissipation rate of high SiCrFe and graphite anodes is actually 0.5 kg/A-year or less, it is common practice to use 1.0 kg/A-year to allow for a safety factor (similar to the utilization factor used in galvanic anode design). For other anode materials, the designer should follow the dissipation information given by the manufacture. The total anode weight (Wt) for the desired life then is Wt ¼ Dr Ireq L

ð69:10Þ

where Wt is the total weight required in kilograms, Dr is the dissipation rate in kg/A-year, Ireq is current requirement in amperes, and L is life in years. The number of anodes required is simply the total weight required divided by the individual anode weight.

1008


F3.5. Groundbed Resistance F3.5.1. Anode-to-Ground Resistance. Since multiple anodes are usually involved, Eq. (69.9) is used to calculate the resistance of the anode bed to earth (for vertical anodes), as discussed in Section E4 for groups of anodes in a galvanic system. The dimensions of the coke column are used for L and d in the equation. F3.5.2. Cable Resistance. Figure 69.4 shows a typical surface point groundbed layout. There are three cable resistances to consider: Negative cable to the pipeline Positive cable to the groundbed Positive cable in the anode groundbed

F3.5.4. Total Circuit Resistance. The total circuit resistance in ohms (Rt) is the sum of the anode to ground (Ra), cable (Rc), and structure-to-ground (Rs) resistances, or R t ¼ Ra þ R c þ R s

ð69:11Þ

where Rc is the total cable resistance, R( ) is the negative cable resistance, R( þ ) is the positive cable resistance, and Rgb is one-half of the resistance of the total length of the anode portion of the groundbed. Cable resistances are given in Table 69.3.

ð69:12Þ

F3.6. Rectifier Selection. Required driving voltage (E) is determined by Ohm’s law: E¼

Since the current flow through the anode portion of the groundbed cable drops as each anode is encountered, the effective resistance of the groundbed cable is usually taken as one-half of its length. Thus the total effective cable resistance in ohms becomes Rc ¼ Rð Þ þ Rð þ Þ þ Rgb

F3.5.3. Structure-to-Ground Resistance. As discussed in Section E, the pipe-to-ground resistance is frequently negligible. Peabody [4, p. 135] provides additional information on calculating structure-to-ground resistance.

Ireq Rt

ð69:13Þ

The designer should allow for some increase in current requirement over the years; 25% is a reasonable figure, but the designer should also rely on experience. It is prudent also to allow for some circuit resistance increase in the future, so the required rectifier voltage is usually multiplied by 1.5. The designer then chooses a commercially available unit meeting the design requirement. F4. Other Groundbed Configurations A distributed anode system, shown in Figure 69.5, consists of anodes spaced at intervals along the structure to be protected. Typical applications include pipelines, sheet pilings, large tank farms, and other complex networks.

FIGURE 69.4. Impressed current cathodic protection using surface point groundbed.

DESIGN OF OTHER CATHODIC PROTECTION SYSTEMS

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FIGURE 69.6. Impressed current cathodic protection using a deep anode groundbed.

The structure being protected affects the design. Stationary facilities such as bulkheads and support piles require different techniques from ship hulls. Morgan [3, Chapter 7] discusses this in detail. FIGURE 69.5. Impressed current cathodic protection using distributed groundbed.

Figure 69.6 shows a deep anode bed. This installation utilizes anodes placed in a coke column vertically in the earth. Typical depths are 60–90 m. The resistance to earth is based on the dimensions of the coke breeze column and may be calculated from Eq. (69.2). Lewis [8] presents detailed information on the design and use of deep anode groundbeds.

G. DESIGN OF OTHER CATHODIC PROTECTION SYSTEMS Not all uses of cathodic protection can be explained in this chapter. The following sections provide references for various applications. G1. Marine Structures Design of cathodic protection for marine structures in both fresh and salt water require special techniques. Galvanic systems usually employ zinc or aluminum alloy anodes. Impressed current systems frequently use high silicon, chromium-bearing iron, platinized niobium, or mixed-metal oxide/titanium anodes.

G2. Heat Exchangers Cathodic protection is often used to overcome corrosion in heat exchanger water boxes. It is especially useful in controlling galvanic corrosion between steel tube sheets and copper or other noble metal tubes such as titanium or stainless steel. Morgan [3, p. 398 ff] and Lane [9] present information on this application. G3. Steel in Concrete Bridge decks, parking structures, and other reinforced concrete structures lend themselves to cathodic protection, particularly when the concrete is contaminated with chlorides from deicing salts or a marine environment. Special techniques are required as described by Rog and Swiat [10], NACE [11, 12] and Morgan [3, p. 242 ff]. Prestressed concrete pipelines occasionally require cathodic protection. Protection must be done carefully to avoid damage to the prestressing wire from hydrogen embrittlement or stress corrosion cracking (SCC). Helpful advice is found in [13]. G4. Water Storage Tanks The wetted surfaces of water storage tank interiors can easily be protected by cathodic protection. Above the water surface,

1010


the tank must be protected by coating. Refer to Bushman and Kroon [14] and Kumar et al. [1] for design details.

several small installations to ensure that all areas are protected.

G5. Above-Ground Storage Tank Bottoms

H3. Cathodic Interference

Cathodic protection is used on both Single- and doublebottom tanks. Many jurisdictions require such protection. Morgan [3, p. 232 ff] provides design details.

This term refers to the stray-current effect that a cathodic protection installation may have on other structures. Interference is seldom a problem with galvanic anode protection because of its inherent small current output and driving voltage. With impressed current, however, straycurrent effects may occur on adjacent structure. This needs to be taken into consideration in the design of the design of the protection. See [15, Chapter 5] for information on this topic.

H. OTHER DESIGN CONCERNS H1. Electrical Continuity In most cases, the structure to be protected must be electrically continuous. This is particularly true for piping, reinforcing steel, bulkheads, foundations piles, and the like. Where electrical continuity does not exist, such as in piping with mechanical joints or among many structures in a complex network, joint bonds, consisting of insulated copper cable, are necessary to ensure that all of the structures to be protected are electrically continuous. There are times where electrical continuity is not required. One example is the installation of galvanic anodes on existing ductile iron pipe where anodes can be installed on individual lengths of pipe.

H4. Attenuation For long pipelines many kilometers in length, the attenuation, or reduction of protection with distance, must be considered. This becomes increasingly important as the pipelineto-electrolyte resistance decreases or as the linear resistance of the pipeline increases. Figure 69.7 shows typical attenuation on a pipeline. Specific information on this aspect of cathodic protection engineering has been presented by Morgan [3, p.201 ff].

I. CRITERIA FOR CATHODIC PROTECTION

H2. Shielding In complex situations such as tank farms, industrial plants, and other large underground and underwater structures, it is important to lay out the anodes so that all structures receive protection. This may require a distributed anode system or

The designer must understand the applicable criterion and the proper test techniques. Various criteria for underground and submerged structures are listed below. The reader must study these criteria and associated testing techniques in detail.

1.000 0.980

Pipe to Soil Potential, V

0.960 0.940 0.920 0.900 0.880 0.860 Location of CP source

0.840 0.820 0.800 0

1

2

3

4

5

6

7

8

9

10 11 12 13 14 15 16

Distance, km

FIGURE 69.7. Typical attenuation on a pipeline.

REFERENCES

See [16] regarding underground and submerged structures and Section 2 of [12] for steel in concrete. I1. Steel and Cast Iron A negative potential of at least 850 mV to a saturated copper–copper sulfate reference electrode (CSE) with the cathodic protection current applied. Voltage drops other than those across the structure to electrolyte boundary must be considered. A negative polarized potential of at least 850 mV with respect to CSE. A minimum of 100 mV of cathodic polarization between the structure and a stable reference electrode. This criterion also applies to steel in concrete. I2. Aluminum A minimum of 100 mV of cathodic polarization between the structure and a stable reference electrode. Precautions must be taken to prevent overprotection of aluminum. I3. Copper A minimum of 100 mV of cathodic polarization between the structure and a stable reference electrode.

REFERENCES 1. A. Kumar, J. B. Bushman, J. H. Fitzgerald, A. E. Brown, and T. M. Kelly, Impressed Current Cathodic Protection Systems Utilizing Ceramic Anodes, U.S. Army Corps of Engineers Construction Engineering Research Laboratories, Champaign, IL, 1990.

1011

2. Basic Course, Text, Appalachian Underground Corrosion Short Course, West Virginia University, Morgantown. WV, 2008. 3. J. H. Morgan, Cathodic Protection, 2nd ed., NACE International, Houston, TX, 1993. 4. A. W. Peabody, Control of Pipeline Corrosion, 2nd ed., NACE International, Houston, TX, 2001. 5. Advanced Course Text, Appalachian Underground Corrosion Short Course, West Virginia University, Morgantown, WV, 2008. 6. H. B. Dwight, “Calculation of Resistance to Ground,” Electric. Eng., 55(12), 1319 (1936). 7. E. D. Sunde, Earth Conduction Effects in Transmission Systems, D. Van Nostrand, New York, 1949. 8. T. H. Lewis, Deep Anode Systems, NACE International, Houston, TX, 2008. 9. R. W. Lane, Control of Scale and Corrosion in Building Water Systems, McGraw-Hill, New York, 1993. 10. J. W. Rog and W. J. Swiat, “Guidelines for Selection of Cathodic Protection Systems for Reinforced Concrete,” paper presented at Corrosion 87, San Francisco, NACE International, Houston, TX, 1987. 11. V. Chaker (Ed.), Corrosion Forms and Control for Infrastructure, NACE International, Houston, TX, 1992. 12. Standard Practice SP02-90-2007, “Impressed Current Cathodic Protection of Reinforcing Steel in Atmospherically Exposed Concrete Structures,” NACE International, Houston, TX, 2007. 13. M. Szeliga (Ed.), Corrosion in Prestressed Concrete: Pipes, Piles and Decks, NACE International, Houston, TX, 1995. 14. J. B. Busman and D. H. Kroon, J. AWWA, Jan. 1984. 15. Intermediate Course Text, Appalachian Underground Corrosion Short Course, West Virginia University, Morgantown, WV, 2004. 16. Standard Practice SP01-69-2007, “Control of External Corrosion on Underground or Submerged Metallic Piping Systems,” Section 6, NACE International, Houston, TX, 2007.

70 STRAY-CURRENT ANALYSIS J. H. FITZGERALD III Grosse Pointe Park, Michigan

A. Introduction B. Detection of stray currents C. Locating source of interference C1. Static interference C2. Dynamic stray currents D. Mitigation of stray-current corrosion D1. Controlling stray currents at source D2. Static stray currents D3. Dynamic stray currents D4. Galvanic anodes D5. Impressed current systems E. Stray current from ac sources References

A. INTRODUCTION The causes and common means of detecting and mitigating stray-current interference effects from direct-current (dc) sources are reviewed in this chapter. Alternating current (ac), which can create potential safety hazards has also been shown to contribute to corrosion of ferrous structures. Alternating ac stray current is discussed in Section E. Stray currents are defined as electrical currents flowing through electrical paths other than the intended paths. Stray currents, or interference currents, are classified as either static or dynamic. Static interference currents are those that maintain constant amplitude and constant paths. Examples of typical sources are railroad signal batteries, high-voltage direct-current (HVDC) *

Adapted with permission from the Appalachian Underground Corrosion Short Course [1, 2].

ground electrodes, and cathodic protection system rectifiers. Dynamic interference currents are those that continually vary in amplitude, magnitude, and electrolytic paths. These currents can be man made (e.g., dc welding equipment, dc railway systems, chloride plants, and aluminum plants) or caused by natural phenomena. Natural sources of dynamic stray currents, called tellurics, are caused by disturbances in the earth’s magnetic field from sun spot activity. Telluric effects may contribute to corrosion and, in addition, can create measurement difficulties and interfere with the ability to assess cathodic protection system performance. The lower the resistivity of the soil, the more severe the effects of stray currents may be, If there is a current flowing in the earth and a potential difference exists between points where a metallic conductor, such as a pipeline or cable, is located, then the conductor will readily acquire a part of the current that is flowing. Thus, metallic pipelines and cables can become conductors of stray currents in the earth environment.

B. DETECTION OF STRAY CURRENTS Static stray currents on a pipeline can be detected by analyzing pipe-to-soil potentials. The graph in Figure 70.1 shows a pipeline with no interference. Figure 70.2 shows a potential plot for a coated pipeline with stray-current interference. Interference may be suspected if: The voltage curve profile shows abnormal variation from previous survey graphs. High negative values are noted remote from any cathodic protection system on the surveyed line. Unusual current flows are measured along the pipeline. Low negative or positive voltages are measured.


1013

1014

STRAY-CURRENT ANALYSIS

FIGURE 70.1. Potential versus distance plot, no interferance.

FIGURE 70.2. Potential plot with interference; coated pipeline.

In practice, if the positive shift does not cause a potential less negative than 0.850 vs. Cu/CuSO4 (free of voltage IR drop caused by current flow through ohmic resistance), then negligible corrosion can be expected. Dynamic stray currents are more easily detected than static ones. Dynamic stray currents are present if the structure-to-soil potential is continually fluctuating while the reference electrode is kept in a stationary position in contact with the soil. These potential changes are the result of current changes at the source of the interference.

C. LOCATING SOURCE OF INTERFERENCE C1. Static Interference The path of current flow in the earth can be tracked to its source by measuring the currents in the earth using two identical portable reference electrodes and a digital voltmeter. By measuring the potential difference between the two electrodes spaced about 8 m (25 ft) apart, as shown in Figure 70.3, the direction of current flow can be determined

LOCATING SOURCE OF INTERFERENCE

1015

FIGURE 70.3. Tracing stray-current flow.

and its path traced. With the voltmeter connected as shown, positive potential readings would indicate current flow toward the pipeline. Negative potential readings would indicate that current is flowing away from the pipeline. If there are test facilities for measuring current flow along the pipeline, a line current survey can be conducted to determine the areas of current pick-up and discharge. This information can also be used to track stray currents toward their source. To confirm the cause of abnormal pipe-to-soil potential readings, the effect of interrupting the suspected current source on a potential survey along the protected pipeline or structure should be assessed. If there is no effect on the potential readings of the interfered structure, the search for the current source must continue until the actual interfering current source is located.

C2. Dynamic Stray Currents After identifying possible sources, such as dc electrical railway systems, mines, or industrial plants, such as aluminum and chlorine, the current flow should be traced along the interfered structure to its source. One method is to observe the current flow at intervals along the structure using millivolt drop test station lead wires to determine the direction of current flow, as illustrated in Figure 70.4. Assume that a situation exists where a single source is causing interference problems. A voltmeter connection, to be used for pipe-to-soil potential measurements, is made between the pipeline (interfered structure) and a reference electrode within the earth current pattern of the source and its load. Observing the fluctuating potential readings at this point alone would not enable one to determine if the readings are being taken at a point where the pipeline is picking up or

FIGURE 70.4. Determining direction of line current flow.

1016


FIGURE 70.5. Typical test setup used to locate point of maximum exposure.

discharging current. If measurements are being taken at a point of current pickup, a negative potential swing would be indicated. A positive swing would indicate a decrease of current flow and a condition of the pipe returning to its steady-state condition. Readings observed at a discharge point swinging in the positive direction would indicate an increase in current leaving the line and a negative swing would indicate a decrease in current discharge. Test locations can be identified as pick-up or discharge points by correlating the open-circuit voltage between the interfered pipeline and the stray-current source and the pipe-to-soil potential of the pipeline. A plot of these data is called a beta plot. Figure 70.5 illustrates a typical experimental setup to obtain a beta plot. A current pickup area is indicated by pipe-to-soil potentials (Vg) that become more negative as pipe-to-negative bus voltages (E1) increase; Figure 70.6(a) shows the beta plot for a pickup area. A current discharge area is indicated by pipe-to-soil potentials (Vg) that become less negative as (pipe-to-negative bus voltages (E1) increase; Figure 70.6(b) shows the beta plot for a discharge area. As the connection shown in Figure 70.5 indicates, a positive value of E1 occurs when current flows from the stray-current source to the pipeline. Conversely, a negative value of E1 indicates a current flow from the pipeline to the stray-current source. The slope of a straight line through the data points is the value of b ¼ DVg/DE. The beta curve (slope) of dynamic stray-current data is the opposite of the geometric slope; the geometric slope would be DE1/DVg. This type of testing requires that pipe-to-soil potential Vg be measured at two or more locations. Many readings should be taken at these locations simultaneously and the meters used must be identical or comparison of the sets of readings will be difficult. In most cases, a dual-channel recorder, such as an X–Y plotter, or a multichannel data logger is used.

If the points plotted form a vertical line, a neutral curve, there is no influence on Vg by the output fluctuations of the current source. The point of maximum exposure to stray currents can be determined from a plot of slopes of beta curves versus distance along the pipeline. The location of maximum discharge area slope (the line closest to the horizontal) is the location of maximum exposure.

D. MITIGATION OF STRAY-CURRENT CORROSION D1. Controlling Stray Currents at Source Groundbed site selection can be used to eliminate or greatly reduce stray currents and the potential gradients that accompany them. Ideally, groundbeds should be installed as far as possible from any foreign structure in the area to minimize the effects of the electric field from the groundbed. In transit systems and any of the other systems involving rail returns, the rails should be installed on wellballasted road beds or on insulated ties or padding with ungrounded substations. Similarly, when dealing with equipment, if isolated positive and negative circuits can be employed, stray-current problems will be minimized because of the high circuit resistance to earth. When welding is done, care should be taken to ensure that the ground connection and the welding electrode are relatively close together and that the electrical path between them is of negligible resistance. D2. Static Stray Currents Mitigation bonds are used to mitigate the effects of straycurrent corrosion on a structure. The purpose of the miti-

MITIGATION OF STRAY-CURRENT CORROSION

1017

FIGURE 70.6. Typical beta curve: (a) pickup area; (b) discharge area.

gation bond is to eliminate current flow from a metallic structure into the earth by providing a metallic return path for the current. This bond allows the stray current flowing from a groundbed to the interfered structure to flow through the structure and back to the protected structure through the bond. The typical current flow when a drainage bond is installed is shown in Figure 70.7. Corrosion will occur only if the current flows from the metal surface into the earth. To size the mitigation bond, the point of maximum current flow between the two affected structures must be located. Typically, this point is situated near the point of pipeline crossing, where the circuit resistances are the lowest, but it can be located some distance away, particularly

with well-coated pipelines at areas of coating failure or damage. A “trial-and-error” method can be used to determine the correct bond resistance and current for a solution to static stray currents. A reference electrode is placed at the point of maximum stray-current exchange to monitor the potentialto-soil of the interfered structure. With the current source operating, a variable resistor is placed between the two pipelines. When the potential-to-soil of the interferred line with the current source operating and the bond installed equals the pontential-to-soil of the interferred line with the current source deactivated and with the bond disconnected, the correct bond resistance is determined. The current flow

1018


FIGURE 70.7. Typical current flow through resistive bond.

through the bond and the resistance of the bond can then be measured. The required bond resistance can also be calculated mathematically using the following equation:

See [1] for a detailed discussion of static stray-current location and mitigation.

D3. Dynamic Stray Currents E0 þ DE=ISC RI Rb ¼ Ib where

Rb ¼ resistance of mitigation bond RI ¼ resistance between structure E0 ¼ open-circuit potential without stray-current source operating DE ¼ change in open-circuit potential caused by stray-current source ISC ¼ current flow from stray-current source Ib ¼ current in mitigation bond

Once the location of maximum exposure is determined and its slope or beta curve plotted, the resistance of the bond can be determined. The required size of the resistance bond is such that its installation will cause the beta curve at the point of maximum exposure to assume a neutral or pick-up slope. Figure 70.8 shows a beta curve at a point of maximum exposure as well as the required mitigation curve. In sizing the mitigation bond, a trial-and-error solution may be possible in relatively simple cases where a single source of stray current is involved. The size of the resistance

FIGURE 70.8. Beta curve—discharge area (mitigation current).

STRAY CURRENT FROM AC SOURCES

bond can be determined by installing a temporary variable resistance and determining when stray-current corrosion has been mitigated. The procedure is similar to that described for static stray currents, except that a mitigation curve, such as that shown in Figure 70.8, must be obtained. Where a more complex interference problem exists, which precludes the use of the trial-and-error method, a mathematical method can be used. In this method, the following equation is used to obtain the required resistance of the mitigation bond: Rb ¼ where

DVg =Ib RI b

Rb ¼ resistance of mitigation bond, W RI ¼ resistance between structures, W Ib ¼ bond current, A b ¼ beta slope DVg ¼ change in pipe-to-soil potential caused by Ib

Using values determined experimentally by field measurements, Rb, the resistance of the bond, is calculated. For example, if DVg =Ib ¼ 0:00169 V=A

b ¼ 0:017

RI ¼ 0:070 W

then Rb ¼ 0:0294 W The bond resistance value, 0.0294 W for this example, can be a simple cable or a combination of cable and variable or fixed resistors. The bond must also be sized to permit the maximum current to flow and remain in the current range for the bond. In a dynamic stray-current situation, the maximum current can be calculated once the maximum open-circuit voltage between the structures is known. This value is usually obtained by measuring the park value over several operating cycles of the stray-current source. For most stray-current sources, the typical cycle is 24 h. For example, if the maximum value of open-circuit voltage E1 (Fig. 70.5) is 12.0 V, the value of the maximum stray current through the bond can be calculated as follows: E1 ¼ Ib ðRI þ Rb Þ

Ib ¼ 120:9 A

See [2] for detailed analysis and calculations. In many instances of stray current where rail transit systems are involved, there may be locations where drainage bonds are required in areas where reversals of potential could occur. Therefore, it is often necessary to install an electrolysis (reverse-current) switch or silicon diode into the circuit to prevent current flow from the substation back onto the pipeline through the bond. The resistance to the forward

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flow of current created by these devices must be included in the sizing of the bond.

D4. Galvanic Anodes Galvanic, or sacrificial, anodes may be used to mitigate straycurrent effects in situations where small current flows or small voltage gradients exist. In effect, a potential gradient produced by the galvanic anode(s) counteracts the interference current, with a resulting net current flow to the interfered-with structure. Because galvanic anodes produce limited voltages, they can overcome only limited straycurrent voltages. In addition, resistance of galvanic anodes to ground increases as the anodes are consumed during their lifetime, leading to reduced current flow from the anode and decreased voltage gradients. Galvanic anodes should be sized to provide a sufficient anticipated life span and should be monitored carefully. Galvanic anode drains are commonly used in lieu of bonds where small drain currents are involved.

D5. Impressed Current Systems When the magnitude of stray currents is beyond the ability of galvanic anodes to counteract, impressed current systems can sometimes be utilized. Impressed current systems have much higher voltage capacities than galvanic anodes and a greater life per kilogram of anode material. Built-in control and monitoring circuits in the impressed current rectifier can be used to adjust the protective current output based on the voltage-to-earth fluctuations of the interfered-with pipeline. The design of galvanic or impressed current mitigation systems is done by the trial-and-error method. Simulated systems are placed in the field and the results measured. Based on the results, a full-scale system can be designed.

E. STRAY CURRENT FROM AC SOURCES Stray current from ac sources is a well-known safety hazard, particularly in the form of ac voltages that can occur on pipelines paralleling or crossing high-voltage ac electrical tower lines. Recent work, however, is leading to the understanding that ac discharge from a pipeline may indeed cause corrosion [3]. Both steady-state and dynamic ac stray current may be found. See [4] for detailed information. Steady-state ac is commonly associated with pipelines laid in close proximity or paralleling high-voltage ac electrical transmission lines. Dangerous voltages can be created on such pipelines. Mitigation of the voltage as well as safety grounding of test stations and above-ground appurtenances to buried equipotential mats are recommended if ac pipe-to-soil potentials exceed 15 V.

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Dynamic ac stray currents may be generated by ac welding, bad building grounds, or ac electrified railroads.

REFERENCES 1. Intermediate Course Text, Appalachian Underground Corrosion Short Course Omit, West Virginia University, Morgantown, WV, 2008, Chapter 5.

2. Advanced Course Text, Appalachian Underground Corrosion Short Course, West Virginia University, Morgantown, WV, 2008, Chapter 4. 3. R. Gummow, “AC Corrosion—A Challenge to Pipeline Integrity,” Mater. Perform. 38(2) (Feb. 1999). 4. NACE: Standard Practice SP01-77-2007, “Mitigation of Alternating Current and Lightning Effects on Metallic Structures and Corrosion Control Systems,” NACE International, Houston, TX, 2007.

71 CORROSION INHIBITORS S. PAPAVINASAM CANMET Materials Technology Laboratory, Hamilton, Ontario, Canada

A. Introduction B. Classification of inhibitors B1. Environmental conditioners (scavengers) B2. Interface inhibitors B2.1. Liquid-phase inhibitors B2.1.1 Anodic inhibitors B2.1.2 Cathodic inhibitors B2.1.3 Mixed inhibitors B2.2. Vapor-phase inhibitors C. Mechanistic aspects of corrosion inhibition C1. Environmental conditioners (scavengers) C2. Anodic inhibitors (passivators) C3. Cathodic inhibitors C4. Mixed inhibitors (adsorption) C5. Vapor-phase inhibitors D. Examples of corrosion inhibitors D1. Inhibitors containing oxygen atom D2. Inhibitors containing nitrogen atom D3. Inhibitors containing sulfur atom D4. Electronically conducting polymers D5. Coordination complexes E. Industrial applications of corrosion inhibitors E1. Petroleum production E2. Internal corrosion of steel pipelines E3. Water E4. Acids E5. Automobile E6. Plants (organic coatings) E7. Miscellaneous F. Other factors in applying inhibitors F1. Application techniques F2. Temperature effects F3. Poisoning F4. Secondary inhibition

F5. Synergism and antagonism F6. Green inhibitors References

Her Majesty the Queen in Right of Canada, as represented by the Minister of Natural Resources, 2010.

Inhibitor selection is based on the metal and the environment. A qualitative classification of inhibitors is presented in

A. INTRODUCTION A corrosion inhibitor is a chemical substance which, when added in small concentrations to an environment, minimizes or prevents corrosion [1]. Corrosion inhibitors are used to protect metals from corrosion, including temporary protection during storage or transport as well as localized protection, required, for example, to prevent corrosion that may result from accumulation of small amounts of an aggressive phase. One example is brine in a nonaggressive phase such as oil. An efficient inhibitor is compatible with the environment, is economical for application, and produces the desired effect when present in small concentrations. Inhibitor efficiency P is given as P¼

w0 w 100 w0

ð71:1Þ

where wo is the corrosion rate in the absence of inhibitor and w is the corrosion rate in the same environment with the inhibitor added.

B. CLASSIFICATION OF INHIBITORS


1021

1022

CORROSION INHIBITORS

FIGURE 71.1. Classification of inhibitors: () Form three-dimensional layers at the interface, so they are classified collectively as interphase inhibitors.

Figure 71.1. Inhibitors can be classified into environmental conditioners and interface inhibitors. B1. Environmental Conditioners (Scavengers) Corrosion can be controlled by removing the corrosive species in the medium. Inhibitors that decrease corrosivity of the medium by scavenging the aggressive substances are called environmental conditioners or scavengers. In nearneutral and alkaline solutions, oxygen reduction is a common cathodic reaction. In such situations, corrosion can be controlled by decreasing the oxygen content using scavengers (e.g., hydrazine [2]). B2. Interface Inhibitors Interface inhibitions control corrosion by forming a film at the metal/environment Interface. Interface inhibitors can be classified into liquid- and vapor-phase inhibitors. B2.1. Liquid-Phase Inhibitors. Liquid-phase inhibitors are classified as anodic, cathodic, or mixed inhibitors, depending on whether they inhibit the anodic, cathodic, or both electrochemical reactions. B2.1.1 Anodic Inhibitors. Anodic inhibitors are usually used in near-neutral solutions where sparingly soluble corrosion products, such as oxides, hydroxides, or salts, are formed. They form, or facilitate the formation of, passivating films that inhibit the anodic metal dissolution reaction. Anodic inhibitors are often called passivating inhibitors. When the concentration of an anodic inhibitor is not sufficient, corrosion may be accelerated, rather than inhibited. The critical concentration above which inhibitors are effective depends on the nature and concentration of the aggressive ions.

B2.1.2 Cathodic Inhibitors. Cathodic inhibitors control corrosion by either decreasing the reduction rate (cathodic poisons) or by precipitating selectively on the cathodic areas (cathodic precipitators). Cathodic poisons, such as sulfides and selenides, are adsorbed on the metal surface, whereas compounds of arsenic, bismuth, and antimony are reduced at the cathode and form a metallic layer. In near-neutral and alkaline solutions, inorganic anions, such as phosphates. silicates, and borates, form protective films that decrease the cathodic reaction rate by limiting the diffusion of oxygen to the metal surface. Cathodic poisons can cause hydrogen blisters and hydrogen embrittlement due to the absorption of hydrogen into steel. This problem may occur in acid solutions, where the reduction reaction is hydrogen evolution, and when the inhibitor poisons, or minimizes, the recombination of hydrogen atoms to gaseous hydrogen molecules. In this situation, the hydrogen, instead of leaving the surface as hydrogen gas, diffuses into steel causing hydrogen damage, such as hydrogen-induced cracking (HIC), hydrogen embrittlement, or sulfide stress cracking. Cathodic precipitators increase the alkalinity at cathodic sites and precipitate insoluble compounds on the metal surface. The most widely used cathodic precipitators are the carbonates of calcium and magnesium. B2.1.3 Mixed Inhibitors. About 80% of inhibitors are organic compounds that cannot be designated specifically as anodic or cathodic and are known as mixed inhibitors. The effectiveness of organic inhibitors is related to the extent to which they adsorb and cover the metal surface. Adsorption depends on the structure of the inhibitor, on the surface charge of the metal, and on the type of electrolyte. Mixed inhibitors protect the metal in three possible ways: physical adsorption, chemisorption, and film formation.

MECHANISTIC ASPECTS OF CORROSION INHIBITION

FIGURE 71.2. Adsorption of negatively charged inhibitor on a positively charged metal surface.

Physical (or electrostatic) adsorption is a result of electrostatic attraction between the inhibitor and the metal surface. When the metal surface is positively charged, adsorption of negatively charged (anionic) inhibitors is facilitated (Fig. 71.2). Positively charged molecules acting in combination with a negatively charged intermediate can inhibit a positively charged metal. Anions, such as halide ions, in solution adsorb on the positively charged metal surface, and organic cations subsequently adsorb on the dipole (Fig. 71.3). Corrosion of iron in sulfuric acid containing chloride ions is inhibited by quaternary ammonium cations through this synergistic effect [3].

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Physically adsorbed inhibitors interact rapidly, but they are also easily removed from the surface. Increase in temperature generally facilitates desorption of physically adsorbed inhibitor molecules. The most effective inhibitors are those that chemically adsorb (chemisorb), a process that involves charge sharing or charge transfer between the inhibitor molecules and the metal surface. Chemisorption takes place more slowly than physical adsorption. As temperature increases, adsorption and inhibition also increase. Chemisorption is specific and is not completely reversible [4]. Adsorbed inhibitor molecules may undergo surface reactions, producing polymeric films. Corrosion protection increases markedly as the films grow from nearly twodimensional adsorbed layers to three-dimensional films up to several hundred angstroms thick. Inhibition is effective only when the films are adherent, are not soluble, and prevent access of the solution to the metal. Protective films may be nonconducting (sometimes called ohmic inhibitors because they increase the resistance of the circuit, thereby inhibiting the corrosion process) or conducting (self-healing films). B2.2. Vapor-Phase Inhibitors. Temporary protection against atmospheric corrosion, particularly in closed environments, can be achieved using vapor-phase inhibitors (VPIs). Substances having low but significant pressure of vapor with inhibiting properties are effective. The VPIs are used by impregnating wrapping paper or by placing them loosely inside a closed container [5]. The slow vaporization of the inhibitor protects against air and moisture. In general, VPIs are more effective for ferrous than nonferrous metals.

C. MECHANISTIC ASPECTS OF CORROSION INHIBITION C1. Environmental Conditioners (Scavengers) In near-neutral solutions, the common cathodic reaction is oxygen reduction: O2 þ 2H2 O þ 4e > 4OH

ð71:2Þ

Scavengers deplete the oxygen by chemical reaction; for example, hydrazine removes oxygen by the following reaction [2]: 5O2 þ 2ðNH2 NH2 Þ > 2H2 O þ 4H þ þ 4NO2

ð71:3Þ

C2. Anodic Inhibitors (Passivators) FIGURE 71.3. (a) Positively charged inhibitor molecule does not interact with positively charged metal surface. (b) Synergistic adsorption of positively charged inhibitor and anion on a positively charged metal surface.

The mechanism of anodic inhibition can be explained using the polarization diagram of an active–passive metal (Fig. 71.4) [6].

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FIGURE 71.4. Polarization diagram of active–passive metal showing dependence of current on concentration of passivation-type inhibitor [7].

In the absence of inhibitors, the metal corrodes in the active state at a rate corresponding to point A in Figure 71.4. As the concentration of inhibitor is increased, the corrosion rate also increases until a critical concentration and a critical corrosion rate (point B, Fig. 71.4) are reached. At the critical concentration, there is a rapid transition of the metal to the passive state, and the corrosion rate is decreased (point C). C3. Cathodic Inhibitors In acid solution, the cathodic reaction is, typically, the reduction of hydrogen ions to hydrogen atoms, which combine forming hydrogen molecules:

Hþ þ e > H

ð71:4Þ

2H > H2

ð71:5Þ

In alkaline solution, the cathodic reaction is typically oxygen reduction [Eq. (71.2)]. Cathodic inhibitors impede reduction reactions. Substances with high overpotential for hydrogen and those that form precipitates at the cathode are effective in acid and alkaline solutions, respectively. The effect of a cathodic inhibitor on the polarization curves is shown in Figure 71.5. In this case, the slope of the anodic polarization curve is

FIGURE 71.5. Polarization curve in the presence of cathodic inhibitor [7].

MECHANISTIC ASPECTS OF CORROSION INHIBITION

unaffected, but the slope of the cathodic polarization curve is changed [7]. C4. Mixed Inhibitors (Adsorption) Adsorption occurs as a result of electrostatic forces between the electric charge on the metal and the ionic charges or dipoles on the inhibitor molecules. The potential at which there is no charge on the metal is known as the zero-charge potential (ZCP) (Table 71.1) [8]. The charge on a metal surface in a given medium can be determined from the corrosion potential (ECORR) and zerocharge potential. When the difference ECORR – ZCP is negative, the metal is negatively charged and adsorption of cations is favored. When ECORR – ZCP is positive, the metal is positively charged and adsorption of anions is favored. The charge on inhibitors depends on the presence of loosely bound electrons, lone pairs of electrons, p-electron clouds, aromatic (e.g., benzene) ring systems, and functional groups containing elements of group V or VI of the periodic table. Most organic inhibitors possess at least one functional group, regarded as the reaction center or anchoring group. The strength of adsorption depends on the charge on this anchoring group [rather on the hetero atom (i.e., atoms other

TABLE 71.1. Metal Ag Al Au Bi Cd Co Cr Cu Fe Ga Hg In Ir Nb Ni Pb Pd Pt Rh Sb Sn Ta Ti Tl Zn a

See [8].

Values of Zero-Charge Potentialsa Zero-Charge Potential, mV (SHE) 440 520 þ 180 390 720 450 450 þ 90 350 690 190 650 40 790 300 620 0 þ 20 20 140 430 850 1050 750 630

1025

than carbon, including nitrogen and sulfur) present in the anchoring group]. The structure of the rest of the molecule influences the charge density on the anchoring group [4]. Water molecules adsorb on the metal surface immersed in an aqueous phase. Organic molecules adsorb by replacing the water molecules: ½Inhibitorsoln þ ½nH2 Oadsorbed > ½Inhibitoradsorbed þ ½nH2 Osoln ð71:6Þ where n is the number of water molecules displaced by one inhibitor molecule. The ability of the inhibitor to replace water molecules depends on the electrostatic interaction between the metal and the inhibitor. On the other hand, the number of water molecules displaced depends on the size and orientation of the inhibitor molecule. Thus, the first interaction between inhibitor and metal surface is nonspecific and involves low activation energy. This process, called “physical adsorption,” is rapid and, in many cases, reversible [9]. Under favorable conditions, the adsorbed molecules involved in chemical interaction (chemisorption) form bonds with the metal surface. Chemisorption is specific and is not reversible. The bonding occurs with electron transfer or sharing between metal and inhibitor. Electron transfer is typical for transition metals having vacant, low-energy electron orbitals. Inhibitors having relatively loosely bound electrons transfer charge easily, The inhibition efficiency of the homologous series of organic substances differing only in the heteroatom is usually in the following sequence: P > S > N > O. An homologous series is given in Figure 71.6. On the other hand, the electronegativity, that is, the ability to attract electrons, increases in the reverse order. Adsorption strength can be deduced from the adsorption isotherm, which shows the equilibrium relationship between concentrations of inhibitors on the surface and in the bulk of the solution. Various adsorption isotherms to characterize inhibitor efficiencies are presented in Table 71.2 [10–13]. To evaluate the nature and strength of adsorption, the experimental data (e.g., corrosion rate) are fitted to the isotherm, and from the best fit, the thermodynamic data for adsorption are evaluated.

FIGURE 71.6. Homologous series of organic molecules (the molecules differ only in the heteroatom.)

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TABLE 71.2. Adsorption Isothermsa Name

Isothermb

Verification Plot

Langmuir Frumkin Bockris– Swinkels

q/(1/q) ¼ bc ½q=ð1 qÞef q ¼ b c q=ð1 qÞn ½q þ nð1 qÞn 1 = nn ¼ c e b =55:4 q ¼ ð1=f Þ ln K c q e2f q ¼ b c

q/(1 q) vs. log c q vs. log c q/(1 q) vs. log c

Temkin Virial Parson

q vs. log c q vs. log (q/c)

a

See [10–13]. q, %P/100, surface coverage; b, DG/2.303RT; DG, free energy of adsorption; R, gas constant; T, temperature; c, bulk inhibitor concentration; n, number of water molecules replaced per inhibitor molecule; f, inhibitor interaction parameter (0, no interaction; þ , attraction; and , repulsion); K, constant; and %P ¼ 1 inhibited corrosion rate/uninhibited corrosion rate.

b

The principle of soft and hard acids and bases (SHAB) has also been applied to explain adsorption and inhibition [14]. The SHAB principle states that hard acids prefer to coordinate with hard bases and that soft acids prefer to coordinate with soft bases. Metal atoms on oxide-free surfaces are considered to be soft acids, which in acid solutions form strong bonds with soft bases, such as sulfur-containing organic inhibitors. By comparison, nitrogen or oxygencontaining organic compounds are considered to be hard bases and may establish weaker bonds with metal surfaces in acid solutions. Whatever may be the mechanism of adsorption, the electron density of the functional groups, polarizability, and electronegativity are important parameters that determine inhibition efficiency.

TABLE 71.3. Some Anchoring (Functional) Groups in Organic Inhibitors Structure

Name

Structure

Name

OH C¼C COC COOH CNC NH2 NH NO2 N¼NN

Hydroxy -Yne Epoxy Carboxy Amine Amino Imino Nitro Triazole

CONH2 SH S S¼O C¼S P¼O P As Se

Amide Thiol Sulfide Sulfoxide Thio Phosphonium Phospho Arsano Seleno

some industrial chemical processes for which the active constituent is not known. Commercial inhibitor packages may contain, in addition to the active ingredients for inhibition, other chemicals, including surfactants, deemulsifiers, carriers (e.g., solvents), and biocides. The active ingredients of organic inhibitors invariably contain one or more functional groups containing one or more heteroatoms, N, O, S, P, or Se (selenium), through which the inhibitors anchor onto the metal surface. Some common anchoring groups are listed in Table 71.3. These groups are attached to a parent chain (backbone), which increases the ability of the inhibitor molecule to cover a large surface area. Common repeating units of the parent chain are methyl and phenyl groups. The backbone may contain additional molecules, or substituent groups, to enhance the electronic bonding strength of the anchoring group on the metal and/or to enhance the surface coverage. The outline of the constitution of an organic inhibitor is presented in Table 71.4.

C5. Vapor-Phase Inhibitors The process of vapor-phase inhibition involves two steps: transport of inhibitor to the metal surface and interaction of inhibitor on the surface. Either a VPI may vaporize in the molecular form or it may first dissociate and then vaporize. Amines vaporize in the undissociated molecular form. Subsequent reaction with water, present as moisture at the surface, dissociates the inhibitor. On the other hand, dicyclohexylamine nitrite dissociates liberating amine and nitrous acid, which deposit on the metal surface [15]. Both in molecular and in dissociated forms VPls adsorb either physically or chemically on the metal surface to inhibit corrosion.

D. EXAMPLES OF CORROSION INHIBITORS Inhibitors used in practice are seldom pure substances but are usually mixtures that may be byproducts, for example, of

D1. Inhibitors Containing Oxygen Atom Benzoic acid and substituted benzoic acids are widely used as corrosion inhibitors [16]. Adsorption and inhibitor efficiencies of benzoic acids depend on the nature of the

TABLE 71.4. Constitution of an Organic Corrosion Inhibitor Anchoring Groupa Binds onto the metal

Backbone

Substituent Groupsa

Bears anchoring and substituent groups Provides surface coverage

Supplements electronic strength and surface coverage

a Anchoring and substituent groups are interchangeable, that is, the substituent group through which the inhibitor anchors onto the metal surface depends on the electron density, charge on the metal, and orientation of the molecule in a particular environment.

EXAMPLES OF CORROSION INHIBITORS

1027

FIGURE 71.7. Variation of inhibitor efficiency as a function of substituents (benzoic acid) (substituents with negative Hammett constants will attract electrons from the anchoring COOH group, thereby decreasing the efficiency) [16]. (Hammett constant is a measure of ability of the substituents to attract or repel electrons.)

substituents. Electron-donating substituents increase inhibition by increasing the electron density of the anchoring group (COOH group); on the other hand, electron-withdrawing substituents decrease inhibition by decreasing the electron density. Percent inhibition as a function of substituents is presented in Figure 71.7.

D2. Inhibitors Containing Nitrogen Atom Benzotriazole (BTA) and its derivatives are effective inhibitors for many metals, especially copper, in a variety of conditions [17]. At low concentrations, BTA is adsorbed slightly on the surface. At sufficiently high solution concentrations, bulk precipitation of the complex on the surface

FIGURE 71.8. Dependence of inhibitor efficiency on time (benzotriazole on copper) [17].

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TABLE 71.5. Concentration of Substituted Thioureas Required to Produce 90% Inhibitiona Inhibitor

Chemical Structure

Thiourea Allyl thiourea N,N-Diethyl thiourea N,N-Diisopropylthiourea Phenyl Thiocarbamide Symdiotolylthiourea

H2NCSNH2 H2NCSNHCH2CH¼CH2 C2H5HNCSNHC2H5 C3H7HNCSNHC3H7 H2NCSNHC6H5 C6H5HNCSNHC6H5 CH3C6H4NHCSNHC6H4CH3

Concentration, (mol/L)

Molecular Weight

0.1 0.1 0.003 0.001 0.001 0.0006 0.0004

76.13 116.19 132.23 160.28 152.21 228.38 256.35

a

See [18].

occurs, inhibiting corrosion. Formation of this complex is a slow process and, as a result, the inhibitor efficiency of BTA increases with time (Fig. 71.8). D3. Inhibitors Containing Sulfur Atom Thiourea and its derivatives are used as corrosion inhibitors for a variety of metals [18]. They are nontoxic and are not an environmental hazard. The variation in the inhibitor efficiencies of various derivatives of thiourea depends on the molecular weight. By using lower concentrations of large molecules, higher inhibitor efficiencies can be obtained (Table 71.5.) D4. Electronically Conducting Polymers In situ polymerization of heterocyclic compounds, such as pyrrole and thiophene (structures in Fig. 71.6) and aniline, produces homogeneous, adhesive films on the metal surface [19]. These films are electronically conducting and have the advantage of tolerance to microdefects and minor scratches. Conductivity can be up to 100 S/cm and can be

varied depending on the extent of oxidation, from semiconductor to metal. Because of these properties, the films repassivate any exposed areas of metal where there are defects in the passive film. Conducting polymers are now used as inhibitors for metal corrosion. D5. Coordination Complexes A variety of chelants provide either corrosion inhibition or corrosion acceleration, depending on the structure and functional groups. The chelants displaying high surface activity and low solubility in solution are effective corrosion inhibitors. If they do not have these characteristics, they stimulate corrosion. The 8-hydroxyquinoline molecule satisfies the structural requirements for surface chelation, but formation of a nonadherent film is a distinct disadvantage (Fig. 71.9). On the other hand, pyrocatechols (Fig. 71.10), forming adherent chelants on the metal surface, are effective inhibitors. Inhibition efficiency can be increased by decreasing the solubility through alkylation (increase in chain length) (Table 71.6) [20].

FIGURE 71.9. 8-Hydroxyquinoline surface chelation (stable chelate complex is formed but is soluble in aqueous medium —no corrosion inhibition [20]. (Copyright NACE International. Reprinted with permission.)

INDUSTRIAL APPLICATIONS OF CORROSION INHIBITORS

FIGURE 71.10. Pyrocatechols (forms insoluble chelate complex with the metal. Efficient corrosion inhibitor; also refer to Table 71.6 [20].

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physical adsorption of hydrocarbons (oils) on these nonpolar tails increases film thickness and the effectiveness of the hydrophobic barrier for corrosion inhibition [21]. Because inhibitors are interfacial in nature, they are active at liquid–liquid and/or liquid–gas interfaces and can lead to emulsification. As a result, foaming is sometimes experienced in the presence of inhibitors. E2. Internal Corrosion of Steel Pipelines

TABLE 71.6. Dependence of Inhibitor Efficiency of Pyrocatechol on Chain Lengtha

Inhibitor

Substituent (R value in Fig. 71.10)

% Inhibition

Pyrocatechol 4-Methylpyrocatechol 4-n-Butylpyrocatechol 4-n-Hexylpyrocatechol

–H –CH3 –(CH2)3CH3 –(CH2)5CH3

14 84 93 96

a

See [20]. Copyright NACE International. Reproduced with permission.

E. INDUSTRIAL APPLICATIONS OF CORROSION INHIBITORS

Gathering pipelines, operating between oil and gas wells and processing plants, have corrosion problems similar to those in refineries and petrochemical plants. The flow regimes of multiphase fluids in pipelines influence the corrosion rate. At high flow rates, flow-induced corrosion and erosion– corrosion may occur, whereas at low flow rates, pitting corrosion is more common. Corrosion is related to the amount and nature of sediments. High-velocity flow tends to sweep sediments out of the pipeline, whereas low velocity allows sediments to settle at the bottom, providing sites for pitting corrosion. Internal corrosion of pipelines is controlled by cleaning the pipeline (pigging) and by adding continuous and or batch inhibitors. E3. Water

E1. Petroleum Production Corrosion in the hydrocarbon industries may be divided into two types, “wet corrosion” and “dry corrosion.” At low temperature (i.e., below the boiling point or dewpoint of water), material corrodes due to the presence of an aqueous phase (wet corrosion). At higher temperature (above the boiling point of water), corrosion occurs in the absence of an aqueous phase (dry corrosion). Wet corrosion is inftuenced by pressure, temperature, and compositions of aqueous, gaseous, and oil phases. In refineries and petrochemical plants, the amount of water is usually small, but the corrosivity is high and is localized at regions where the aqueous phase contacts the metal. The water may contain dissolved hydrogen sulfide (H2S), carbon dioxide (CO2), and chloride ions (CI ). Corrosion may occur even when the produced water content is as low as 0.1%, or corrosion activity may not begin until after several years of production. Refineries and petrochemical industries employ a variety of film-forming inhibitors to control wet corrosion. Most of the inhibitors are long-chain nitrogenous organic materials, including amines and amides. Water-soluble and water soluble–oil dispersible inhibitors are continuously injected, or oil-soluble and oil soluble–water dispersible inhibitors (batch inhibitors) are intermittently applied to control corrosion. Film-forming inhibitors anchor to the metal through their polar group. The nonpolar tail protrudes out vertically. The

Potable water is frequently saturated with dissolved oxygen and is corrosive unless a protective film, or deposit, is formed. Cathodic inhibitors, such as calcium carbonate, silicates, polyphosphates, and zinc salts, are used to control potable water corrosion. Water is used in cooling systems In many industries. In a recirculating system, evaporation is the chief source of cooling. As evaporation proceeds, the dissolved mineral salt content increases. Cooling systems may consist of several dissimilar metals and nonmetals. Metals picked up from one part of the system can be deposited elsewhere, producing galvanic corrosion. Corrosion is controlled by anodic (passivating) inhibitors, including nitrate and chromate, as well as by cathodic (e.g., zinc salt) inhibitors. Organic inhibitors (e.g., benzotriazole) are sometimes used as secondary inhibitors, especially when excessive corrosion of copper occurs [22]. E4. Acids Acids are widely used in pickling, in cleaning of oil refinery equipment and heat exchangers, and in oil well acidizing. Mixed inhibitors are widely used to control acid corrosion. E5. Automobile Inhibitors are used in an automobile for two reasons: (1) to reduce the corrosivity of fluid systems (internal corrosion)

1030


and (2) to protect the metal surfaces exposed to the atmosphere (external corrosion). Internal corrosion is influenced by the coolants, flow, aeration, temperature, pressure, water impurities and corrosion products, operating conditions, and maintenance of the system. Some common inhibitors dissolved in antifreeze are nitrites, nitrates, phosphates, silicates, arsenates, and chromates (anodic inhibitors); amines, benzoates, mercaptans, and organic phosphates (mixed inhibitors); and polar or emulsifiable oils (film formers) [23]. Atmospheres to which automobiles are exposed contain moist air, wet SO2 gas (forming sulfuric acid in the presence of moist air), and deicing salt (NaCl and CaCl2). To control external corrosion, the rust-proofing formulations that are used contain grease, wax resin, and resin emulsion, along with metalloorganic and asphaltic compounds. Typical inhibitors used in rust-proofing applications are fatty acids, phosphonates, sulfonates, and carboxylates. E6. Paints (Organic Coatings) Finely divided inhibiting pigments are frequently incorporated in primers. These polar compounds displace water and orient themselves in such a way that the hydrophobic ends face the environment, thereby augmenting the bonding of the coatings on the surface. Red lead (Pb3O4) is commonly used in paints on iron. It deters formation of local cells and helps preserve the physical properties of the paints. Other inhibitors used in paints are lead azelate, calcium plumbate, and lead suboxide [24]. E7. Miscellaneous Inhibitors are used to control corrosion in boiler waters, fuel oil tanks, hot chloride dye baths, refrigeration brines, and reinforcing steel in concrete, and they are also used to protect artifacts.

F. OTHER FACTORS IN APPLYING INHIBITORS Some factors to be considered in applying inhibitors are discussed in the following paragraphs. F1. Application Techniques A reliable method should be applied for inhibitor application. A frequent cause of ineffective inhibition is loss of the inhibitor before it either contacts the metal surface or changes the environment to the extent required. Even the best inhibitor will fail if not applied properly. If the inhibitor is continuously applied in a multiphase system, it should partition into the corrosive phase, usually

the aqueous phase. This partitioning is especially important when using water-soluble, oil-dispersible inhibitors. In batch treatment, the frequency of treatment depends on the film persistency. It is important that the corrosion rates are measured frequently to ensure that a safe level of inhibition is maintained. It is also important that the inhibitor contacts the entire metal surface and forms a continuous persistent film. When using volatile inhibitors, care must be taken in packaging to prevent the loss of inhibitor to the outside atmosphere. Inhibitors are added to the primers used in paint coatings. When moisture contacts the paint, some inhibitor is leached from the primer to protect the metal. The inhibitor should be incorporated in such a way that it protects the areas where potential corrosion can take place and not leach completely from the primer during the service life. F2. Temperature Effects Organic molecules decompose at elevated temperatures. In general, film-forming inhibitors that depend on physical adsorption become less effective at elevated temperatures, so that larger treatment dosages may be required to maintain protective films. Chemisorption, on the other hand, increases with temperature due to the strengthening of chemical bonds. As a result, inhibitor efficiency increases with temperature up to the temperature at which decomposition of the inhibitor occurs. F3. Poisoning Inhibitors for hydrogen damage should reduce not only the corrosion rate but also the rate of absorption and permeation of hydrogen into the steel. For example, the corrosion rate of steel in sulfuric acid is decreased [25], while hydrogen flux through a steel membrane is increased by adding thiourea [26]. Although thiourea inhibits corrosion, it poisons the hydrogen recombination reaction, so that much of the hydrogen produced at the steel surface enters the steel, causing hydrogen damage, rather than recombining with other hydrogen atoms to form bubbles of hydrogen that escape from the system. F4. Secondary Inhibition The nature of the inhibitor initially present in acid solutions may change with time as a consequence of chemical or electrochemical reactions. Inhibition due to the reaction products is called secondary inhibition. Depending on the effectiveness of the reaction products, secondary inhibition may be higher or lower than primary inhibition. For example, diphenyl sulfoxide undergoes electrochemical reaction at the metal surface to produce diphenyl sulfide, which is more effective than the primary compound [27]. On the contrary,

OTHER FACTORS IN APPLYING INHIBITORS

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FIGURE 71.11. Synergistic effect of mixing formaldehyde and furfuralimine [28].

the reduction of thiourea and its alkyl (e.g., methyl, ethyl) derivatives gives rise to HS , which accelerates corrosion. F5. Synergism and Antagonism In the presence of two or more adsorbed species, lateral interaction between inhibitor molecules can significantly affect inhibitor performance. If the interaction is attractive, a synergistic effect arises, that is, the degree of inhibition in the presence of both inhibitors is higher than the sum of the individual effects. For example, because of this synergistic effect, the inhibition efficiency of a mixture of formaldehyde and furfuralimine is higher compared to the inhibition

efficiency when these inhibitors are used separately (Fig. 71.11). On the other hand, when narcotine and thiourea are used as mixed inhibitors, there is an antagonistic effect and a decrease in inhibitor efficiency compared to that which exists when these inhibitors are used separately (Fig. 71.12) [28]. F6. Green Inhibitors Environmental concerns worldwide are increasing and are likely to influence the choice of corrosion inhibitors in the future. Environmental requirements are still being developed, but some elements have been established.

FIGURE 71.12. Antagonistic effect of mixing narcotine and thiourea [28].

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TABLE 71.7. Toxicity of Corrosion Inhibitorsa Compound Propargyl alcohol Hexynol Cinnamaldehyde Formaldehyde Dodecylpyridinium bromide Naphthylmethylquinolinium chloride Nonylphenol–ethylene oxide surfactants

LD50, (mg/kg) 55 34 2200 800 320 644 1310

a

See [30].

The biodegradation, or biological oxygen demand (BOD), should be at least 60%, and inhibitors should be nontoxic [29]. The BOD is a measure of how long the inhibitor will persist in the environment. Toxicity is measured as LC50 or EC50 [30]. LC50 is the concentration of the inhibitor needed to kill 50% of the total population of the test species. The results are quoted as milligrams of chemical per liter of fluid (or LD50, milligrams per kilogram) for exposure times of 24 and 48 h. The EC50 is the effective concentration of inhibitor to adversely affect 50% of the population. In general, EC50 values are lower than LC50 values because the former are the concentrations required to damage the species in some way without killing it. Some chemicals are excellent inhibitors but are quite toxic and readily adsorbed through the skin. Toxicity of some inhibitors is presented in Table 71.7. There is a growing demand for corrosion inhibitors that are less toxic and biodegradable compared to current formulations. Green inhibitors displaying substantially improved environmental properties will be the inhibitors most widely used in the future.

REFERENCES 1. O. L. Riggs, Jr., in C. C. Nathan (Ed.), Corrosion Inhibitors, NACE, Houston, TX, 1973, p. 11. 2. M. G. Noack, Mater. Perform., 21(3), 26 (1982). 3. A, Frignani, G. Trabanelli, F. Zucchi, and M. Zucchini, in Proceedings of 5th European Symposium of Corrosion Inhibitors, Ferrara, Italy 1980, p. 1185. 4. V. S. Sastri, Corrosion Inhibitors: Principles and Applications, Wiley, New York, 1998, p. 39. 5. S. A. Levin, S. A. Gintzbergy, I. S. Dinner, and V. N. Kuchinsky, in Proceedings of Second European Symposium on Corrosion Inhibitors, Ferrara, Italy 1965, p. 765.

6. M. G. Fontana, Corrosion Engineering, 3rd ed., McGraw-Hill, New York, 1986, pp. 445–481. 7. N. Hackerman and E. S. Snavely, in Corrosion Basics, L. S. V. Delinder (Ed.), NACE, Houston, TX, 1984, pp. 127–146. 8. S. Trasatti, J. Electroanal. Chem. Interf. Electrochem., 33, 351 (1971). 9. W. Lorenz, Z. Phys. Chem., 219, 421 (1962); 224, 145 (1963); and 244, 65 (1970). 10. Z. S. Smialowska and G. Wieczorek, Corros. Sci., 11, 843 (1971). 11. S. Trasatti, J. Electroanal. Chem., 53, 335 (1974). 12. J. O. M. Bockris and D. A. J. Swinkels, J. Electrochem. Soc., 111, 736 (1964). 13. I. Langmuir, J. Am. Chem. Soc., 39, 1848 (1947). 14. R. G. Pearson, J. Am. Chem. Soc., 85, 3533 (1963); Science, 151, 172 (1966). 15. I. L. Rosenfeld, V. P. Persiantseva, and P. B. Terentief, Corrosion, 20, 222t (1964). 16. A. Akiyama and K. Nobe, J. Electrochem. Soc., 117, 999 (1970). 17. P. G. Fox, G. Lewis, and P. J. Boden, Corros. Sci., 19, 457 (1979). 18. I. Singh, Corrosion, 49, 473 (1993). 19. T. A. Skotheim (Ed.), Handbook of Conducting Polymers, Vols. 1 and 2, Marcel Dekker, New York, 1986. 20. A, Weisstuch, D. A. Carter, and C. C. Nathan, Mater. Perform., 10(4), 11, (1971). 21. C. C. Nathan (Ed.), Corrosion Inhibitors, NACE, Houston, TX, 1973, p. 45. 22. S. Matsuda and H. H. Uhlig, J. Electrochem. Soc., 111, 156 (1964). 23. Snow and Ice Control with Chemical and Abrasives, Bulletin 152, Highway Research Board, Washington, DC, 1960. 24. T. Rossel, Werkstoffe Korrosion, 20, 854 (1969). 25. R. W. Revie and H. H. Uhlig, Corrosion and Corrosion Control, 4th ed., Wiley, Hoboken, NJ, 2008, p. 312. 26. G. Trabenelli and F. Zucchi, Rev. Coat. Corros., 1, 97 (1972). 27. G. Trabanelli, F. Zucci, G. L. Zucchini, and V, Carassiti, Electrochim. Met., 2, 463 (1967). 28. I. N. Putilova, S. A. Balezin, and V. P. Barannik, Metallic Corrosion, Inhibitors, Pergamon, New York, 1960, pp. 17–24. 29. R. L. Martin, B. A. Alink, T. G. Braga, A. J. McMahon, and R. Weare, “Environmentally Acceptable Water Soluble Corrosion Inhibitors,” Paper No. 36, CORROSION/95, NACE, Houston, TX, 1995. 30. W. W. Frenier, “Development and Testing of a Low-Toxicity Acid Corrosion Inhibitor for Industrial Cleaning Applications,” Paper No. 152, CORROSION/96, NACE, Houston, TX, 1996.

72 COMPUTER TECHNOLOGY FOR CORROSION ASSESSMENT AND CONTROL S. SRINIVASAN Advanced Solutions—Americas, Honeywell International, Inc, Houston, Texas

A. Introduction B. Computer-based corrosion problem solving: classification and background C. Types of computer programs for corrosion control D. Computer tools: evolutionary development trends E. Internet and world wide web: new frontiers in computing E1. Internet resources for corrosion engineer F. Computers in corrosion control: critical evaluation G. Summary References

A. INTRODUCTION Information and knowledge are the foundations on which all technology has been built. Computers have become ubiquitous, key components of automation and optimized problem solving in most domains of science and engineering, and corrosion is no exception. Computer-based information systems and computer models of corrosion processes have revolutionized our approach to problem solving, information access, and knowledge processing in significant ways. It would not be superfluous to say that in the current-day environment every aspect of corrosion data analyses, modeling, and production is managed through computer-based tools and technologies. This chapter provides an overview of the types of computer tools utilized for solving corrosion-related problems, data storage, and data analyses. An introduction to computerbased corrosion problem solving is followed by a description of types of computer programs employed in the domain of

corrosion, with an overview of early systems leading to current-day models for thermodynamic analyses, corrosion prediction, and material selection. Brief descriptions of different types of computer applications for corrosion, including expert systems, neural networks, and object-oriented software systems, is also included for purposes of both completeness and relevance. A description of currently available computer tools for modeling corrosion and cracking problems, selection of materials/equipment, as well as corrosion management, monitoring, and control is also provided. Aspects relevant to the role of emerging technologies and the availability of computer-based tools for corrosion analyses from Web/Internet-based systems to cloud computing (data shared on Internet servers) are discussed. B. COMPUTER-BASED CORROSION PROBLEM SOLVING: CLASSIFICATION AND BACKGROUND Using computer tools to model and represent corrosion processes is a challenging task since characterizing corrosion processes requires a fundamental understanding of principles underlying multiple disciplines, from electrochemistry, thermodynamics, and fluid mechanics to material science and engineering. The complexity of characterizing corrosionrelated tasks has necessitated use of computer tools in corrosion science and engineering, from thermodynamic and phase behavior modeling to data acquisition and analysis. Computers, in the current-day environment, are an intrinsic part of both data representation and automated problem solving. In this context, computer-based corrosion problemsolving systems may be classified as:


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COMPUTER TECHNOLOGY FOR CORROSION ASSESSMENT AND CONTROL

Systems for modeling corrosion/cracking processes Material selection and equipment specification programs Computer-based corrosion monitoring systems Computer-based systems for control of corrosion testing equipment Databases and hypertext systems Internet-based databases and software programs A large number of early programs in corrosion were billed as expert systems, primarily because the programs typically attempted to capture human expertise in corrosion [1], and these programs represented research-based development efforts normally lacking rigorous software engineering foundations necessary for commercial distribution. Most of these programs were developed using software platforms called shells [2] that supported easy implementation of heuristic rules (rules of thumb) and representation of common concepts of reasoning. Table 72.1 lists a few of the early, wellknown computer programs in corrosion developed in the late 1980s and early 1990s [3–8]. It is interesting to note that none of these early systems were implemented in commonly used programming languages (such as Cþ þ , Fortran, etc.) or current programming languages/environments (JAVA, C#, .Net, etc.), and many were implemented by corrosion/ materials specialists with little or no formal training in software development [9]. Most current-day applications for corrosion may be broken into different categories as follows: Programs for prediction of corrosion for different industrial applications, such as oil and gas production, transmission, and refining Applications for characterizing thermodynamics and phase behavior as well as system speciation for a range of process/plant applications (broadly termed ionic modeling and process modeling tools) Applications for selection of materials for corrosion resistance as well as resistance to stress corrosion cracking Applications for risk-based inspection and risk-based integrity characterization

The single most popular application of computers in corrosion stems from programs built for modeling and predicting corrosion for different applications related to oil and gas. Such programs (currently commercially available) include PredictÒ [9, 10] and NORSOCÒ [11] as well as programs such as HydroCorr from Shell and Cassandra from BP, all of which focus on predicting corrosion in CO2/H2S oil/gas production environments. Numerous other systems for corrosion prediction include Predict-Pipe [12] for internal corrosion direct assessment (ICDA) automation for drygas pipelines and Corrosion Analyzer [13] for characterizing thermodynamics/phase behavior/speciation in corroding systems. Some early systems, such as Auscor [14], provided assessment of stress corrosion cracking in austenitic stainless steels, and legacy programs for general corrosion prediction included LipuCor [15] and Cormed [16]. A more detailed description of some of the current-day corrosion prediction models are provided in a subsequent section in this chapter. Several corrosion-related applications and databases are available to assist with material selection and equipment specifications. The CORIS system [17] is a legacy expert system database that integrates corrosion problem-solving expertise with a comprehensive database on corrosion, materials, and thermodynamic properties of corrosive media. Other programs for material evaluation and selection, some of which have had wide commercial application, that provide an automated basis for materials selection include Ò SOCRATES [18, 19] for oil and gas production, Predict SW [20, 21] for refinery sour water corrosion prediction and material selection and Selmatel [22] for selection of materials exposed to elevated temperatures in refinery furnace tubes. Early applications of computers to corrosion (in the 1990s) focused on development of databases integrated with material evaluation heuristics for corrosion and materials data. The Materials Technology Institute [1] in the late 1980s and early 1990s developed a series of systems called ChemCor relevant to evaluation of materials applicable to different segments of the chemical processing industries. CORSUR [23] was another large database application on the corrosion behavior of metals and nonmetals in over

TABLE 72.1. Early, Legacy Software (Expert) Systems in Corrosion Name

Application

Country of Development

Shell Used

Achilles Aurora Auscor ChemCor DIASCC KISS Prime Suscept

Diagnosis and prediction of localized corrosion Prediction of localized corrosion in austenitic steels Prediction of corrosion of austenitic stainless steels Materials selection for hazardous chemical service Evaluation of cracking in stainless steels Material selection Materials selection for chemical process industry Evaluation of SSC in steels

United Kingdom Finland United Kingdom United States Japan Germany Belgium United States

Spices Level 5 Savior KES/Level 5 OPS83 Nexpert Object KEE PCPLUS

TYPES OF COMPUTER PROGRAMS FOR CORROSION CONTROL

700 chemical environments. However, both ChemCor and Corsur have become obsolete, since MTI has not pursued further development and commercialization of these programs beyond the 1990s.

C. TYPES OF COMPUTER PROGRAMS FOR CORROSION CONTROL Different types of software technologies have found application in corrosion engineering and science. These programs can be broadly classified as follows: Conventional software systems (structured software systems) Artificial intelligence and expert systems applications Object-oriented software systems (includes current-day component-based development frameworks such as .Net and Java) Neural networks Hybrid systems that utilize one or more of the above technologies Most early computer programs and legacy systems (meaning older systems whose technology is currently obsolete) developed in the 1970s and 1980s used highlevel programming languages such as Basic, Fortran, and Pascal. Typically, such systems provided a front end for data modeling and analysis that was computation intensive but provided little or no support for representation of heuristic concepts and manipulation of symbolic information. The advent of expert systems in the late 1980s drastically changed the direction of computer programs for corrosion applications. Expert systems or knowledge-based

systems, defined as intelligent computer programs that use expert knowledge to attain high levels of performance in narrow problem domains [24], became quite common and prominent. The term expert in the expert systems implied the narrow specialization and competence of a human expert embodied in the system. Expert systems typically had a specific structure that distinguishes them from conventional computer programs in that the domain knowledge (knowledge base) is usually separated from the reasoning process (inference engine), as shown in Figure 72.1. The knowledge base housed the expertise that is embodied in an expert system. The inference engine controlled the manner or the logical path used by the expert system to access the information (facts and rules) in the knowledge base to make decisions. Knowledge representation in an expert system referred to the scheme used to represent a given piece of information in the system. Decision making in an expert system was represented through interconnected rules (rules of thumb that correlate facts, data, and expertise) and objects (descriptions of different system components). Expert systems represent by far the most widely used segment of computer applications in corrosion. Numerous expert systems have been developed to address different aspects of corrosion [1] and include systems for applications in cathodic protection [25], assessment of stress cracking in light water nuclear reactors [26], and prediction of localized corrosion of stainless steels [27]. With the advent of object-oriented systems and the concept of reusable entities that contain both the data and the procedures relevant to an object, the need for separation of knowledge and reasoning as found in expert systems vanished. Object-oriented systems provided a framework for modeling through simulation of behavior of real-life systems. An object in this context is any entity characterized by data (properties) and procedures (methods) for manipulating that data. In an object-centric view of the world,

Knowledge Acquisition Subsystem

Knowledge Base Knowledge Engineer

User Interface Inference Engine

Domain Expert

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Explanation Subsystem

FIGURE 72.1. Expert systems: a simple schematic.

End User

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System Template

Object Library

Object 1

Domain Database

Object 2

Param1, Param2.........Paramn

Consultation object1

Object n

Data Normalization

Consultation object2

System Characterization States

FIGURE 72.2. Structural framework of a generic development system.

computation is behavior simulation [28]. Objects represented as computational abstractions are simulated complete with the characteristics of the simulated objects. Once the abstraction has been adequately characterized, other objects can be derived from the abstraction thus created and such derived objects may be allowed to inherit the properties already specified for the abstraction. Most current-day programs and applications are built using object-oriented frameworks (such as .Net) or languages (C þþ , C#, or Java). Genera [29] is an object-oriented framework for development of problem-solving systems in corrosion and materials and has been used to develop numerous problemsolving systems. The Genera system uses the following framework for implementation, a schematic representation for which is given in Figure 72.2: Material and corrosion systems can be represented as objects whose states can be defined in terms of critical parameters and interrelationships between these parameters. States of different properties of objects are encapsulated within the objects. The object-based framework represents a class of problems in terms of abstractions of commonalities between different domains. For example, an environment could be an environment relevant to refineries or pipelines or a producing well. An abstract view of the environment is that of an object which can be characterized in terms of certain critical variables or operating parameters, such as pressure, temperature, and composition. The variables themselves might vary. However, the environment object can still be represented as a function of the states of certain variables and their interrelationships. Neural networks represent another exciting application of computing technology to corrosion. Evolving from neurobiological insights, neural network technology gives a computer system the capability to actually learn from input data. Artificial neural networks, as neural networks are commonly

referred to (because neural networks by themselves mean biological neural systems found in humans and other carbonbased organic life forms), have provided solutions to problems normally requiring human observation and thought processes. Some real-world applications in corrosion include corrosion data modeling and prediction [30], corrosion data reduction [31], and electrochemical impedance spectroscopy data analysis [32]. Neural networks perform computation in a manner quite different from that used by conventional computers, where a single central processing unit sequentially dictates every segment of activity. Neural networks are built from a large number of very simple processing elements that individually deal with pieces of a big problem. A processing element (PE) simply multiplies an input by a set of weights, and a nonlinearity function transforms the result into an output value. The principles of computation at the PE level are deceptively simple. The power of neural computation comes from the massive interconnection among the PEs which share the load of the overall processing task and from the adaptive nature of the parameters (weights) that interconnect the PEs. Normally, a neural network can have several layers of PEs. The most basic and commonly used neural network architecture is the multilayer perceptron (MLP). Figure 72.3 illustrates a simple MLP. The circles are the PEs arranged in layers. The top row is the input layer, the middle row (there can be many of these) is the hidden layer, and the bottom row is the output layer. The lines represent weighted connections (i.e., a scaling factor) between PEs. The performance of an MLP is measured in terms of a desired signal and an error criterion. The output of the network is compared with a desired response to produce an error. An algorithm called back-propagation is used to adjust the weights a small amount at a time in a way that reduces the error. The network is trained by repeating this process many

COMPUTER TOOLS: EVOLUTIONARY DEVELOPMENT TRENDS

INPUT LAYER HIDDEN LAYER (there may be several hidden layers)

OUTPUT LAYER

Input Vector

Input Layer

Hidden Layer

Output Layer Output Vector

FIGURE 72.3. Schematics for a multilayered perceptron of an artificial neural network.

times. The goal of the training is to reach an optimal solution based on the performance measurement. Current-day computer programs for corrosion assessment and control utilize one or more of the computing technologies described above. For instance, the Strategy [33] programs for evaluation of hydrogen-induced cracking in steels combine principles of expert systems and object-oriented programming, and the work done by Silverman et al. [30] integrates expert systems with neural networks. The next section provides a brief description of specific computer programs utilized for corrosion modeling, prediction, monitoring, and control.

D. COMPUTER TOOLS: EVOLUTIONARY DEVELOPMENT TRENDS This section will focus on examples of current-day commercial applications in corrosion that have find widespread use and application. It is important to note that numerous applications of computers and software technology have been reported over the last two decades. Just in the area of expert system development, over 57 systems were

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reported by two surveys [34] conducted by the Materials Technology Institute and the European Federation of Corrosion (EFC) in the late 1990s. If databases and other types of applications are included, the number of computer programs/software systems for corrosion prediction and control extend well into the hundreds. However, it would be instructive to know about systems that define the evolving trends and those that have contributed to the continuous growth of computing technology to corrosion applications. Table 72.2 provides a listing of different types of representative computer applications in corrosion currently available in the industry. In this section, specific examples of current computer systems from different types of computing applications in corrosion are provided. These include: Cormed, a program for corrosion prediction and assessment SOCRATES, a software system for material selection for oil and gas applications Predict-SW, a program for corrosion prediction and material selection for refinery sour water applications (an example case study demonstrating integration with a process modeling system) and a framework for realtime risk-based integrity and corrosion Analysis Programs for online corrosion monitoring and control Corsur, a database system for evaluation of metals and nonmetals An application of neural networks to corrosion data reduction The Cormed corrosion prediction model [16] developed by Elf (a French oil and gas company) predicts the probability of corrosion in wells. It uses a detailed analysis of field experience on CO2 corrosion as well as data from other published computer models of corrosion. The model uses only a limited number of factors, such as CO2, in situ pH, or bicarbonates for assessing the risk of tubing perforation in production wells. Other comprehensive corrosion prediction computer models include Predict [9, 10], LipuCor [15], and the USL [35] model, systems that examine a larger number of parameters in predicting corrosive damage. The SOCRATES system, developed by InterCorr International, and now by Honeywell International, provides access to the material selection decisions and decision logic of a domain expert and significant experience on utilization and selection of corrosion-resistant alloy (CRA) materials and steels. The primary basis for the system data comes from two large Joint Industry Projects (JIP) conducted by Honeywell for CRA data development, which is embodied in the rules built into SOCRATES. SOCRATES also incorporates data sources such as published literature on as well as field experience related to oil and gas field service.

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TABLE 72.2. Representative Sample of Current-Day Software Programs for Corrosion Application

Type of System

Corrosion analyzer: Faciliates simulation of aqueous and nonaqueous multicomponent systems from a standpoint of thermodynamic behavior and speciation CP diagnostic: Troubleshooting and diagnosis of sacrificial anode and impressed current cathodic protection systems EIS data extrapolation: Uses neural networks to train on electrochemical impedance spectroscopy data for extrapolation Filter debris analysis (FDA) expert system: condition monitoring of aircrafts Genera: Generic problem-solving framework for characterizing corrosion and materials problems LipuCor: Prediction of corrosion in oil and gas systems NORSOK: Prediction of CO2 corrosion in oil and gas production and transmission environments Predict: Prediction of corrosion in CO2/H2S multiphase oil and gas production/transmission environments Predict-Pipe: Automation of internal corrosion direct assessment for gas transmission pipelines Predict-SW: Corrosion prediction and material selection for refinery sour water applications SOCRATES: Selection of materials for oil and gas production service

Object-oriented system implemented in Cþþ integrated with relational databases

Strategy: Programs for evaluation of cracking in steels used in pipelines and refineries USL corrosion model: Program for prediction of corrosion in gas condensate wells

The flow of data in the SOCRATES system is shown in Figure 72.4. Even though a specific order of decision making is indicated by the hierarchy shown in Figure 72.4, the objectoriented implementation ensures that data specification does not have to proceed in any specific sequence. At level 1, the initial set of applicable materials is obtained by determining the application for which the CRA is to be selected. If the application is not known, all classes of materials known to the system become part of the solution set. The system characterizes the environment for severity of general corrosion, localized corrosion, stress corrosion cracking, and other relevant mechanisms using common environmental parameters such as H2S and CO2, bicarbonates, pH, operating temperatures, chlorides, aeration/sulfur, water-to-gas ratio/water cut and gas-to-oil ratio as well as metallurgical parameters. Current-day applications have evolved toward availability of real-time, online corrosion prediction and material selection, as opposed to the conventional offline approach. A framework for real-time availability of applications is shown in Figure 72.5. An online, real-time framework ensures that critical decisions related to corrosion problems and upset conditions may be made before the issues becomes a problem.

Shell-based expert system with a database on CP data Artificial neural network (ANN) application

Visual Basic-based interface and knowledge base Object-oriented system implemented in Cþþ Implemented as a conventional structured program Spreadsheet utilizing numerical data relationships Object-oriented programming C#/.Net integrated with a numerical data model Object-oriented programming Cþþ integrated with a numerical data model Object-oriented programming C#/.Net integrated with a numerical data model Object-oriented expert system implemented in C#/.Net interfacing with a relational database on materials and compositions Implemented in Cþþ and integrated with databases in Microsoft Access Implemented as an expert system in Visual Basic

The framework shown in Figure 72.5 is possible only through availability of integrated corrosion prediction (PredictÒ -SW) and process modeling (Unisim Design) applications [36]. PredictÒ -SW ensures that process equipment is properly specified to deliver desired product throughput and specifications consistent with demands of the operating parameters and potential for material/corrosion damage. The tool helps evaluate effects of feed changes, upsets and equipment downtime on process safety, reliability, and profitability. It also improves plant control, operability, and safety using dynamic simulation (sensitivity analyses) of planned and existing plants. PredictÒ -SW quantifies corrosion rates for the most common construction materials of construction utilized in refineries allowing planning for complex and difficult-toquantify refinery processes. PredictÒ -SW utilizes a quantitative engineering database and decision support model to predict corrosion in alkaline sour water systems as a function of critical environmental parameters, such as NH4HS concentration, H2S partial pressure, temperature, hydrocarbon content, and chemical treatments integrated with characterization of flow regimes and wall shear stress.


Application type Strength Hardness Mechanical Condition

H2S co2 Bicarbonates Temperature Chlorides Sulfur Gas to Oil Ratio Water to Gas Ratio

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Application rules Application Requirements Application Specific CRAs from Database Rules for Env. Severity

Environment Characterization

Evaluation of Steels and Inhibitors

Evaluation for Stress Corrosion Cracking

SCC Rules

Alloys Database Max. Temp Env. Severity

Evaluation for Localized Corrosion

Pitting rules

SSC and Special Requirements Cost Analysis

Literature Database search Recommended Set of CRAs

MR0175 / ISO 15156-2 Rules

FIGURE 72.4. Flow of data in the SOCRATES system for selection of corrosion-resistant alloys.

PredictÒ -SW is the resultant product of this extensive research and can:

. .

.

. . . . .

Predict corrosion rates for a wide range of applicable conditions for 14 commonly used materials, including carbon steel, stainless, Monel, and C-276. Perform flow modeling, compute wall shear stress, and analyze flow regimes in multiphase flow. Evaluate parametric effects with a sensitivity analysis tool. Account for effects of light and heavy hydrocarbons. Correlate flow effects with corrosion rate based on extensive laboratory data and flow modeling. Access data used to support system decision making and analyses.

.

. .

Quantify, characterize, and analyze sour water systems helping to prevent unscheduled shutdowns. Facilitate multipoint analysis and data sharing using Microsoft Excel, Microsoft Word, or email. Quantify corrosion in alkaline sour water systems as a function of NH4HS concentration, velocity (shear stress), H2S partial pressure, temperature, and various additional parametric variables. Seamlessly share data with other analysis tools and modeling systems. Share data with OPC-compatible control systems such as Honeywell’s ExperionÒ Process Knowledge System (PKS), as shown in Figure 72.6. The figure details functional interactions in real-time process modeling using UniSimÒ Design and process intelligence

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COMPUTER TECHNOLOGY FOR CORROSION ASSESSMENT AND CONTROL Quantified Corrosion

Prediction Models

Real Time

Operations

Management

Inspection Work

Inspection Limits (Design)

Real-time

Execute Inspections Evaluate Results Report Nonconformance

Continuous Automated Field Inspection / Collection

Deviation Legend Traditional Risk Based Inspection Program

Currently Available Corrosion Technology

Future Vision for Automated Corrosion

FIGURE 72.5. Framework for real-time deployment of corrosion applications for risk-based management.

FIGURE 72.6. Automated asset integrity framework reliability and risk management in process plants.


.

provided by PredictÒ -SW as a means to enhanced asset integrity and value retention. Easily customize PredictÒ -SW (through Honeywell) and deploy the system on a Web-based framework.

One of the most prominent neural network applications in corrosion facilitates electrochemical impedance spectroscopy (EIS) data extrapolation [32]. EIS is reported to be by far the most popular technique for obtaining instantaneous corrosion rate information. However, one limitation of EIS data is the inability of conventional techniques for validation and extension of the data frequency range. Here, artificial neural networks have been used to extrapolate impedance data to a wider range of frequencies. Another important application of neural networks is a system that identifies the problems associated with the inherent variability of corrosion data and attempts to develop solutions to these problems [31]. These solutions are considered to be general advances in neural network technology for use with noisy or poor-quality data. In essence, they depend on the production of two further output parameters in addition to the corrosion rate. The first of these is the variance and is essentially the same as the conventional statistical parameter. The second has been termed the confidence and is an indicator of the reliability of the prediction of the neural network for a given set of input parameters. he Ecorr (acronym for engineering corrosion) [36] computer-aided learning (CAL) package represents a novel application of computing to corrosion education. This system uses advances in multimedia technology to promote interactive, computer-aided learning. In the Ecorr approach, the student is presented with a series of problems to solve by reference to supporting information. The program itself is built as a string of books wherein information is presented in terms of “objects” such as text, photographs, video, or sound objects on a page. The system consists of a theory base, case study modules, a Control PC

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glossary, and a control center to coordinate navigation and information flow between the user and different modules. The theory base contains basic corrosion information that a student can use while working with a specific case study module. The organization of the system facilitates learning by users with different levels of corrosion proficiency due to the inherent ease of navigation. Currently the program is being designed to provide 12 case study modules at two levels of proficiency (basic/level 1 and advanced/level 2). Real-time computer-based control systems represent an important development in the application of computing to corrosion monitoring and control. Figure 72.7 shows the structure of an automated constant extension rate system (ACert) used for conducting corrosion and cracking evaluation tests. This system facilitates test initiation and conduct and data collection and analysis from a software system integrated with a data acquisition control system and data measurement devices. Corrosion tests can be run with minimal human intervention, and such systems promote corrosion evaluation of difficult-to-simulate environments in an automated manner. One of the problems in conducting conventional laboratory tests stems from the fact that the time required to evaluate cracking can take months of exposure. Slow-strain-rate (SSR) tests (also referred to as constant-extension-rate tests) may be used to mechanically accelerate testing with the application of a slow constant extension rate. Such constant-extension-rate tests promote speeding up the crack initiation process and promote evaluation of material to stress corrosion cracking in a matter of hours or days. The Acert Constant Extension Rate Tester is a computer-controlled closed-loop feedback system facilitating testing over a wide range of testing speeds without the need of cumbersome gear changes. It also has an anti-backlash feature for application of cyclic loads and stresses. An Acert system can be programmed to conduct a variety of tests, including Control Instrumentation

CERT Software

Motor Indexer

Data Acquisition Hardware

Data Acquisition Panel

Stepper Motor

End User Drive

Parametric Inputs (Load, Strain, T, P)

FIGURE 72.7. Schematic representation of automated constant-extension-rate tester.

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constant-extension-rate tests, cyclic slow-strain-rate tests, fracture mechanics tests, and crack growth tests. Hitherto, computing and its applications in corrosion typically dealt with individual or group-based development efforts. The advent of the Internet and the World Wide Web has fundamentally altered computing, communication, and human interaction on a global basis. Now, we have access to a new medium where we are not bound by barriers of differing media in sharing information. Given its ability to combine the best of different media such as printed information, sound, video, and computer files of different types, the Internet has truly revolutionized the manner in which corrosion engineers and scientists across the globe access, process, and share data and knowledge. The next section looks at the impact of the Internet on corrosion computing and the new directions for applications of computers in corrosion.

in corrosion operated through the corrosion server at the Corrosion and Protection Centre at the University of Manchester at e-mail: [email protected]. Blogs and discussion threads online evolved from the early news groups, which were the equivalent of a world wide bulletin board (or a class room) where questions and messages could be posted, comments made, discussions monitored, and answers given. There are currently countless blogs, discussions threads, and websites dedicated to corrosion/materials information Two examples of websites providing dedicated corrosion content are corrosionsource.com and corrosiondoctors.org. The Internet has altered most conventional methods of information access in corrosion science and engineering. Already, significant technical resources are available online, including: .

E. INTERNET AND WORLD WIDE WEB: NEW FRONTIERS IN COMPUTING Since 1996, there has been significant growth in the use of the Internet in sharing and exchanging technical information [37–39]. The growth in terms of utilization of the Internet has been so compelling that the Internet has become the default medium for people to access information, knowledge, and applications [40]. The resources available to the materials/corrosion engineer/scientist fall into different categories and span a wide range and type of relevant topics, from simple descriptions of fundamental mechanisms of corrosion and cracking to technical databases focused on specialized material problems to problem-solving software tools that can be accessed and used on the Internet.

.

.

E1. Internet Resources for Corrosion Engineer The Net offers a variety of resources to promote exchange of information and knowledge in the area of corrosion and materials. The resources may be classified into three broad categories, depending on the mechanism of information exchange: E-mail-based resources Blogs and discussion threads Technical resource websites Electronic mail (e-mail) offers a convenient mechanism to exchange information and computer files. A variation to using e-mail is the notion of a technical mailing list. A mailing list represents a group of people communicating on a specific topic using a simple mail “reflector” program to distribute e-mail messages to various members that have subscribed to the mailing list. A typical example is the CORROS-L [38] list, which is a forum for people interested

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Different Types of Technical Databases. This includes corrosion data on general, localized, galvanic, intergranular and erosion corrosion and cracking data relevant to stress corrosion cracking (SCC), Sulfide stress cracking (SSC), Hydrogen embrittlement cracking (HEC), and so on, as well as materials data on steels, stainless steels, corrosion-resistant alloys, plastics, composites, and so on, all available online Corrosion and Materials Problem-Solving Software Tools. It is now possible to run different software tools for corrosion problem solving from the Internet on a global basis. These programs provide access to state-ofthe-art tools without any of the overheads associated with software procurement and distribution. Archival Laboratory Testing Data and Reports. There is already a significant trend toward providing corrosion scientists and engineers access to archival data and reports online. Several websites provide well-organized indices to data and reports on specific topics, such as general and localized corrosion, different types of steels and elastomers, and different application segments, such as oil and gas, power generation, nuclear industry, and so on. Access to Conference Proceedings. The Internet has significantly altered methods used by people to access archival literature information. It is now possible to organize and provide access to valuable literature data online and thereby obviate the conventional difficulties experienced in disseminating current technology and information on a worldwide basis.

F. COMPUTERS IN CORROSION CONTROL: CRITICAL EVALUATION The last two decades have witnessed explosive growth in the utilization of computers to generate, analyze, and interpret

REFERENCES

corrosion data and knowledge. This growth in computing tools has facilitated both efficient data access/dissemination and the ability to perform tasks hitherto impossible for computer systems and typically performed only by human experts. For instance, expert systems for materials selection now are capable of integrating several complex tasks, including corrosion and cracking evaluation on a large database of materials, application of intelligent rules to determine the best choices and providing the end user with the ability to perform on-the-fly, and context-generated sensitivity analyses. Neural networks for analyzing EIS data provide the ability to learn from the data trends and generate extrapolated impedance data in frequency ranges never seen heretofore. Object-oriented software systems have created reusable software components that can be integrated into multiple problem-solving modules for enhanced operational and functional efficacy. The Internet has not only generated the ability to provide access to an interconnected world of data and information but has also offered a new paradigm for information/knowledge access, exchange, and update. Because of the ability of the Web to seamlessly integrate different media as well as create links without the constraint of a specific order (as in the case of conventional media, such as books), the Web offers a new mechanism to share and disseminate information. The impact of this change in paradigm has resulted in the creation of a “virtual” interconnected world wherein we are able to recreate every efficient aspect of conventional life and use concepts such as the virtual database (an interlinked hyperdatabase that can be seamlessly traversed), the virtual classroom (interactive computer-based learning systems on the Net), the virtual resource room (where one can access software programs and resources), and the virtual conference, a global meeting place for exchange of ideas and information. The applications and evolution of computer software and hardware in corrosion have ushered new paradigms and approaches that will require an altered perspective with respect to problem-solving and technology utilization. Going forward, it is clear that most tasks requiring human effort in the current environment will be replaced by smarter and efficient computer software/hardware systems. In the not-too-distant future, it is conceivable that we will use advanced computing hardware/software to run corrosion tests, collect data, analyze the information, generate reports, and transmit them to all required receiving entities, typically smart computer programs that will make material design and procurement decisions in real time. We are already at the cusp of technology where such immense complexity and knowledge handling (too complex for humans oftentimes) has become reality and the rapidity with which computing technology is evolving has become a reflection of the human spirit opening, doors to a world of endless possibilities in computing.

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G. SUMMARY Numerous computer applications have been developed over the last two decades that have significantly altered our approach to corrosion modeling, prediction, data acquisition, processing, and analysis. Every aspect of evolution in computing technology, including databases, expert systems, object-oriented software systems, and neural networks, has been applied to corrosion science and engineering. A representative sample of such systems has been described and an overview of the relevance of such technology has been provided. The growth potential for computer software and hardware technologies and the Internet given the rapid progress seen in the last 10 years is mind-boggling. However, we can surely predict a future wherein computer tools will play a significant role in addressing most aspects of corrosion problem solving requiring human participation today. This technological growth will come at a cost where we will have to strive to continue to create a higher level of functionality for the displaced human effort. While current technology still requires human intervention at many significant levels, a day in the not-too-distant future appears probable when technological progress is accompanied by the need to create a meaningful role for the human mind, which for all its physical shortcomings vis-a-vis a computer-dominated environment is still eons ahead in evolution when it comes to creativity and intelligence. After all, did we not conceive and create these intelligent systems?

REFERENCES 1. P. R. Roberge, “A Critical Review of Expert Systems for Corrosion Prevention and Control,” Paper No. 368, Corrosion/94, Baltimore, MD, 1994. 2. J. P. Ignizio, Introduction to Expert Systems: The Development and Implementation of Rule-Based Expert Systems, McGrawHill, New York, 1991. 3. The Achilles Club Project—Expert System on Corrosion and Corrosion Control, National Physical Laboratory (NPL), Teddington, Middlesex, UK, Oct. 1986. 4. J. G. Hines and A. Basden, “Experience with the Use of Computers to Handle Corrosion Knowledge,” Br. Corros. J., 26(3), 151–156 (1986). 5. A. Jadot and L. Lancus,“ESCORT: Expert Software for Corrosion Technology,” Dissertation, K. V. Leuven University, Belgium, 1985. 6. G. M. Ugiansky et al., “The NACE-NBS corrosion data program,” Computers in Corrosion Control, J. Fu, R. Heidersbach, and R. Erbar (Eds.), NACE, Houston, TX, 1988. 7. C. S. Fang et al., “Computer Model of a Gas Condensate Well Containing Carbon Dioxide,” Paper No. 465, Corrosion/89, NACE, Houston, TX, 1989.

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8. S. Srinivasan and R. D. Kane, “Expert Systems for Material Selection in Corrosive Environments,” Paper No. 564, Corrosion/90, NACE, Houston, TX, 1990. 9. S. Srinivasan and R. D. Kane, “Prediction of Corrosivity of CO2/H2S Production Environments,” Paper No. 11, Corrosion/ 96, NACE International, Denver, CO, Mar. 1996. 10. S. Srinivasan and R. D. Kane, “Critical Issues in the Application and Evaluation of a Corrosion Prediction Model for Oil & Gas Systems,” Corrosion/2003, Paper No. 640, San Diego, Mar. 2003. 11. R. Nyborg and P. Andersson,“Implementation of CO2 Prediction Models in a Three Phase Fluid Flow Model,” Corrosion/ 2000, Orlando, FL, Mar. 26–31, 2000. 12. V. Lagad, S. Srinivasan, and R. D. Kane,“Facilitating Internal Corrosion Direct Assessment Using Advanced Flow and Corrosion Prediction Models,” Paper No. 08131, Corrosion/ 08, New Orleans, LA, Mar. 2008. 13. A. Anderko and R. D. Young, “Model for Corrosion of Carbon Steel in Lithium Bromide Absorption Refrigeration Systems,” Corros. J., 56(5), 543–555, May 2000. 14. C. P. Sturrock et al., “The AUSCOR Program—An Expert System Both Precocious and Obsolete,” Paper No. 375, Corrosion/94, NACE, Houston, TX, 1994. 15. Y. M. Gunaltun, “Combining Research and Field Data for Corrosion Rate Prediction,” Paper No. 27, Corrosion/96, NACE, Houston, TX, 1996. 16. J. L. Crolet and M. R. Bonis, “Prediction of the Risks of CO2 Corrosion in Oil and Gas Wells,” SPE Production Eng., 6(4), 449 (1991). 17. A. Hatzinosios, “CORIS: A Computer-Based Information System for Corrosion and Corrosion Protection,” Paper No. 373, Corrosion/94, NACE, Houston, TX, 1994. 18. S. Srinivasan, R. D. Kane, and J. W. Skogsberg, “Automated Material Selection and Equipment Specification System for Oil and Gas Production: Concept, Development and Implementation,” Paper No. 134, Corrosion/2003, San Diego, Mar. 2003. 19. S. Srinivasan, Lagad V., and Kane R. D., “Automatic evaluation and selection of corrosion resistant materials for oil and gas production and transmission”, Corrosion/05, Paper No. 05054, Houston, TX, 2005. 20. M. S. Cayard, W. G. Giesbrecht, R. J. Horvath, R. D. Kane, and V. V. Lagad, “Prediction of Ammonium Bisulfide Corrosion and Validation with Refinery Plant Experience,” Paper No. 06577, Corrosion/2006, NACE, Houston, TX, 2006. 21. R. J. Horvath, M. S. Cayard, and R. D. Kane, “Prediction and Assessment of Ammonium Bisulfide Corrosion Under Refinery Sour Water Service Conditions,” Paper No. 06576, Corrosion/2006, NACE, Houston, TX, 2006. 22. A. Valdes et al., “An Expert System for the Selection of Materials Exposed to Elevated Temperatures,” Corrosion/92, Paper No. 275, NACE, Houston, TX, 1992. 23. B. Mashayekhi et al., “Corrosion Data Survey: The Next Generation,” Paper No. 604, Corrosion/97, NACE International, New Orleans, LA, 1997.

24. F. Hayes-Roth et al., Building Expert Systems, Addison Wesley, Reading, MA, 1983, pp. 1–37. 25. V. L. Van Blaricum et al., “Cathodic Protection Diagnostic Expert System,” Paper No. 384, Corrosion/94, NACE, Houston, TX, 1994. 26. M. Urquidi-Macdonald et al., “Object Oriented Expert System to Infer Intergranular Stress Corrosion Cracking in Light Water Nuclear Reactors,” Paper No. 340, Corrosion/90, NACE, Houston, TX, 1990. 27. T. J. Hakkarainen, “A Model for Prediction of Possibility of Localized Corrosion Attack of Stainless Steels,” Paper No. 363, Corrosion/96, NACE, Houston, TX, 1996. 28. I. Pohl, Object-Oriented Programming Using Cþþ , Benjamin Cummings, Redwood City, CA, 1993. 29. S. Srinivasan, “A Generalized Framework for Solving Materials and Corrosion Problems,” Paper No. 368, Corrosion/96, NACE, Houston, TX, 1996. 30. D. C. Silverman and E. M. Rosen, “Corrosion Prediction from Polarization Scans Using an Artificial Neural Network Integrated with an Expert System,” Paper No. 264, Corrosion/92, NACE, Houston, TX, 1992. 31. I. S. Helliwell et al., “Neural Networks for Corrosion Data Reduction,” Paper No. 379, Corrosion/96, NACE, Houston, TX, 1996. 32. M. Urquidi-MacDonald and P. C. Egan, “Validation and Extrapolation of Electrochemical Impedance Spectroscopy Data Analysis,” Corrosion Reviews: Special Issue on Applications of Computers in Corrosion, Guest Ed. P. R. Roberge, Vol. XV, Nos. 1–2, 1997. 33. R. D. Kane and S. Srinivasan “Reliability Assessment of Wet H2S Refinery and Pipeline Equipments: A Knowledge-Based Systems Approach,” in Serviceability of Petroleum, Process and Power Equipment, D. Bagnoli, M. Prager, and D. M. Schlader, PVP Vol. 239, ASME, New York, 1992. 34. P. R. Roberge, “Expert Systems for Corrosion Prevention and Control,” Corrosion Reviews: Special Issue on Applications of Computers in Corrosion, Guest Ed. P. R. Roberge, Vol. XV, Nos. 1–2, 1997. 35. C. D. Adams et al., “Methods of Prediction of Tubing Life for Gas Condensate Wells Containing CO2,” 23rd OTC, Houston, TX, 1991. 36. S. E. Faidi et al., “A Computer Aided Learning Package for Teaching Corrosion to Engineering Students,” Paper No. 377, Corrosion/96, NACE, Houston, TX, 1996. 37. P. Brenden Kahoe, Zen and the Art of the Internet, published onlineatvariouswebsites,www.gutenberg.org/ebooks/34,1996. 38. R. Cottis et al., “The Corrosion Information Server—Experience of the First Year,” Paper No. 376, Corrosion/96, Denver, CO, 1996. 39. S. Srinivasan and R. D. Kane, “Materials and Corrosion Resources on the Internet,” 2nd NACE Latin American Corrosion Congress, Rio de Janeiro, Brasil, Sept. 1996. 40. IDC report on Internet Access Statistics, International Data Corporation (IDC), available: www.idcresearch.com, Apr. 2009.

PART V TESTING FOR CORROSION RESISTANCE

73 PRINCIPLES OF ACCELERATED CORROSION TESTING D. L. JORDAN Ford Motor Company, Dearborn, Michigan

A. Introduction B. Purposes and pitfalls of testing C. Development, selection, and execution of meaningful accelerated corrosion tests C1. Types of information C2. Physical and experimental considerations C3. Statistical considerations D. Concluding remarks References

A. INTRODUCTION Few topics in science and engineering incite as much spirited debate as do the selection, execution, and interpretation of accelerated corrosion tests for raw materials and consumer products. The ongoing dialogue spurred by LaQue’s seminal 1952 paper [1] and the subsequent discussion [2] on salt fog testing has continued for a half-century [3–13] and has been expanded to include the plethora of accelerated corrosion tests with “improved correlation” that have been developed over the decades [14–20]. Despite the huge amount of corrosion research and improved understanding that have been accomplished, one of LaQue’s oft-quoted observations still rings true: “. . . sooner or later one’s judgement may be questioned if continued use of an unreliable test in place of the exercise of good judgement in some other way should lead to a series of poor decisions” [1, p.78]. The fact remains that engineers and technologists need rapid methods to determine the corrosion characteristics and behaviors that take a long time to mature in the natural service environment. The challenge is to develop and exercise good judgment in the selection, conduct, and interpretation of accelerated corrosion tests.

The purpose of this chapter is to encourage the reader to study and understand specific in-service environmental and exposure conditions relevant to particular situations and how those conditions affect various aspects of the overall corrosion behavior of materials in a given application. Only then can we hope to derive benefits from an accelerated corrosion test that includes all relevant stressors and degradation mechanisms. Rather than provide a glossary of popular accelerated corrosion test methods, many of which are discussed in other chapters of this handbook, this chapter introduces some of the many considerations necessary to understand why an accelerated corrosion test may be used for a particular set of conditions. Clearly, the definition of “particular set of conditions” can vary dramatically from case to case and will involve unique definitions or interpretations of the terms test material, environment, corrosive species, protective films, and others, but the principles remain the same. In keeping with the theme of this handbook, “corrosion” is defined as “the degradation of a material by its environment.” The word “degrade” in this context implies that corrosion is a process (i.e., a means of changing from one state to another by way of one or more steps), not a property (i.e., a fixed characteristic that defines a specific state of material). Finally, “accelerated corrosion test” is intended to represent any of a number of cyclic cabinet tests, immersion tests, or electrochemical measurements, all of which are covered in detail by standards organizations [21] and elsewhere in this handbook.

B. PURPOSES AND PITFALLS OF TESTING The primary goal of accelerated corrosion testing is to obtain desired information about the corrosion behavior of a


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material or material system in a particular application in a short period of time. Accelerated corrosion tests are conducted with a variety of purposes, including: Quality control: Determination of batch-to-batch variability. Quality control: Comparison with a reference standard or other acceptance criteria. Study corrosion mechanisms: Probe total system with known factors and measure results. Material selection: Determine the performance ranking of competing materials in a given environment. Development of new corrosion-resistant products: Determine effects of one or more changes to the corrosion protection design, including new materials of construction and corrosion inhibitors. Prediction of service life in specific environments: Test design characteristics and determine safety factors visa`-vis results of comparable in-service performance characterizations. Simulate atmospheric load on a material or object: Avoid time and expense of long-term exposure when specific individual stressors or combinations of stressors are of interest. Simulate specific degradation mechanisms: Determine if hypothesized changes to the engineering design had the expected effect on the degradation mechanism, including the effect on the test environment. Stimulate specific degradation mechanisms: Intentionally introduce or intensify specific stressors and combinations of conditions expected to affect the performance of the engineered system. In all cases, the intent is to obtain approximate or directional data quickly when more exact data take an unacceptably long period of time to develop. It is clear that we must be prepared to provide an accurate description of the operable in-service corrosion mechanisms before we can aspire to reproduce those mechanisms and control their rates masterfully in a meaningful accelerated corrosion test. If that can be accomplished, we may be confident that the accelerated corrosion test provides information that is consistent with the outcome of exposure to some or all of the actual service conditions and is worthy of being used as a tool for material selection and decision making. The difficult task is to find ways to increase the rate of the relevant corrosion processes without introducing extraneous mechanisms or altering (including completely neglecting) the specific degradation mechanisms that occur in the service environment of interest. The “environment” in this context includes not only the natural exposure environment but also all parts of the natural and engineered system that

may directly or indirectly affect the corrosion process in any way. Accelerated corrosion tests that have no proven relationship to the service environment of interest are often irrelevant and can be misleading. With that in mind, there is more to accelerated corrosion testing than just subjecting a material to a standard test and reporting the results. For example, one should not expect information that is immediately applciable to a sulfate exposure to be provided by an accelerated corrosion test that features chloride. A relevant accelerated corrosion test may result from the careful analysis of the corrosion process in the expected service environment and the development and testing of a hypothesis regarding the behavior of the material in a test sequence that contains the appropriate stressors. These stressors include not only elements of the exposure environment itself and all parts of the engineered system but also items that may be inadvertently overlooked [e.g., metallic and nonmetallic fasteners, sheared (work-hardened) edges, identification labels, shipping oils, or any of a number of inservice contaminants].

C. DEVELOPMENT, SELECTION, AND EXECUTION OF MEANINGFUL ACCELERATED CORROSION TESTS C1. Types of Information There are numerous opportunities to measure meaningful quantitative and qualitative parameters that provide a glimpse into the corrosion behavior of the engineered material system at any given point during the corrosion process. The corrosion process that occurs during an accelerated corrosion test may provide reasonable estimates of some, all, or none of the measurable parameters provided by any given in-service environments, which may include polluted and nonpolluted versions of the following: Indoor atmospheres Outdoor atmospheres Total immersion in chemicals, processing fluids, and various waters Intermittent immersion in chemicals, processing fluids, and various waters Practically infinite combinations of the above, with various contributions of temperature, pressure, soluble gases, agitation, and mechanical inputs It is important to decide which corrosion behavior characteristics are of interest for a typical application and test with those in mind. The time to decide is long before conducting the accelerated corrosion test, not during or after.

DEVELOPMENT, SELECTION, AND EXECUTION OF MEANINGFUL ACCELERATED CORROSION TESTS

Typical measurable parameters in accelerated corrosion tests include: Visual appearance of the raw material or consumer product at the end of the test Time to develop the first visual sign of undesirable corrosion products Weight loss Weight gain Corrosion rate (in mm/y, mpy, gmd, or A/m2) Pit depth Pit density Concentration of metal ions in solution Dissolved oxygen consumption in local or bulk electrolyte pH change of local or bulk electrolyte Mechanical or physical properties (e.g., tensile strength, electrical conductivity, toughness) Electrochemical parameters (e.g., corrosion potential, current density, Flade potential) The selection of experimental factors or conditions to include in an accelerated corrosion test depends on what is known about the stressors and corrosion mechanisms in the service environment and what decisions are to be made based on the test results. In addition, the initial, intermediate, and “final” characteristics of corrosion products have a tremendous effect on the corrosion rate in many applications and should be considered to be part of the everchanging local exposure environment during the course of the test [22, 23]. Specific in-service conditions, including geometric configuration, orientation of the corroding surface with respect to rainfall or process fluid flow, soiling by atmospheric fallout and wandering wildlife, proximity to emissions of other corroding or otherwise reactive material systems, exposure to mechanical damage by way of external loading or internal stresses, and differential sun loading, all may seem to be outside the scope of accelerated corrosion testing but could hold the key to the establishment of a meaningful and informative test upon which knowledgeable corrosion-resistant material selection decisions may be made. While it may not be immediately practicable to reproduce certain naturally occurring conditions in the laboratory, there is no reason to ignore their existence while testing and especially during data analysis. Similarly, the interaction of mechanical and environmental stresses during some in-service exposures may require a specific sequence of events during testing. For example, exposure of a specimen to a corrosive environment while under tensile loading (resulting in a yawning crack) will give different results than such an exposure in the nonloaded condition due to the differing abilities of the corrosive

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environment to freely enter, react in, and exit the crack. The exposure conditions within the crack are different in the two loading conditions, so it is reasonable that the corrosion behavior within the crack will differ accordingly. As another example, metals with naturally occurring passive films may be unfairly evaluated when nonrepresentative specimen preparation inadvertently compromises the passive film prior to accelerated corrosion testing. If the in-service conditions are well defined and understood, acceleration of degradation may be realized by exposing thoughtfully prepared specimens to purposefully manipulated combinations of temperature, type and concentration of corrosive species, or time of wetness. On the other hand, incompletely understood service conditions can lead to the selection of inappropriate exposures and acceleration factors that may reduce total testing time but may introduce irrelevant or even misleading results. For example, increasing temperature often results in decreased oxygen solubility in aqueous electrolytes, which in turn leads to unexpectedly low corrosion rates when oxygen reduction is the primary cathodic reaction [4]. Carefully collected and approriately analyzed field data from a range of typical in-service conditions provide the best input to the development of a relevant accelerated corrosion test for future generations of raw materials or consumer products for similar service. There is no substitute for knowing how your product performs when in the hands of your customer. C2. Physical and Experimental Considerations When using an accelerated corrosion test to evaluate modifications to existing corrosion-resistant raw materials or consumer products, it may be tempting to expect an A : B comparison that is always relevant to the service condition to result. The reality is that accelerated corrosion tests with demonstrable and nontrivial correlation with in-service corrosion performance are rare [9, 16, 17], usually due to an incomplete understanding of the in-service degradation mechanisms. The difficulty here is that, for example, an accelerated corrosion test that simply increases temperature and aggressive ion concentration should not be expected, a priori, to highlight differences in coating adhesion, diffusivity of specific anions though a coating, charge transfer at a modified surface, or any of a number of other corrosion-resistant product development strategies that can improve corrosion performance but are much better characterized by propertyspecific tests or measurements that are not normally characterized as “corrosion” tests. Relying on one of the many standard accelerated corrosion tests to reveal improved behavior resulting from a strategy change that is not specifically probed by the standard test can result in the unfair elimination of promising ideas for improved in-service

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performance or the unwitting acceptance of improvement ideas that look good in the test but ultimately disappoint in the field. Know what specific material properties or characteristics are being tested or measured and why. The selection of aggressive ions for accelerated corrosion tests may seem to be a straightforward task. Consider the introduction of the chloride ion to a corrosion test environment. Sodium chloride is an inexpensive, widely studied, and well-accepted source of the frequently destructive chloride ion in aqueous corrosion testing solutions. In addition, sodium chloride is pervasive in daily life and is encountered in the form of road deicers, seawater, and various process fluids that contribute to undesired corrosion events in automotive, aerospace, construction, buried pipeline, and other industries. It seems natural that sodium chloride should be part of many accelerated corrosion tests. Magnesium chloride and calcium chloride are also used a road deicers, and both are more hygroscopic than sodium chloride. Consequently, the time of wetness for a product exposed to the more hygroscopic road deicers will be higher than that of a product exposed to only sodium chloride. In turn, the in-service corrosion rate for materials exposed to the more hygroscopic road deicers may be expected to be higher, all other considerations equal. Accordingly, a continuously wet accelerated corrosion test would not be expected to distinguish the differences in hygroscopicity, time of wetness, or corrosion rate and would fail to detect the true inservice differences in corrosion rates. Not all sources of chloride ions give the same rank order in service as they do in accelerated corrosion tests. The same may be said for other components of accelerated corrosion tests. In accelerated corrosion tests, only certain aggressors are delivered. Some critical aggressors (e.g., thermal cycling to expand and contract surface cracks or to compromise otherwise adherent and protective oxide films) may be key components of in-service conditions but may not be represented at all in an accelerated corrosion test. A limited subset of aggressive ions or species, some of which participate in the initial corrosion reactions and some of which participate in subsequent reactions, and the possible absence of other necessary but not sufficient reactants will allow the painting of only an incomplete picture. Furthermore, some products of initial reactions themselves act as reactants in subsequent reactions [22, 23]. Development of nonrepresentative corrosion products can propagate a series of errors and lead to poor decisions. The presence of a chemical or biological species in the exposure environment or in the resultant soluble and insoluble products of the initial corrosion reactions does not imply that the given species actually participates in the overall reaction. A given species may be an innocent bystander or, at worst, a possible “inhibitor” to the acceleration of the reactions of interest if present in inappropriate concentrations or introduced at an inopportune moment in the overall degra-

dation process (e.g., while conditions are such that a previously yawning crack is held closed). Another practical example of this is where calcium was detected in corrosion products in an automotive body application and a calcium compound was introduced into a developmental accelerated corrosion test. The undissolved portion of the calcium compound clogged nozzles in the test chamber and actually decreased the corrosion rate of materials in the test. The clogging issue was overcome and the intended corrosivity was restored [18], but the exact role of the calcium ions in the test has not yet been established. All is not lost. The point is not to discourage the reader from using accelerated corrosion tests to study and understand corrosion behavior of engineering material and environmental systems but rather to urge careful consideration of how the particular stressors characteristic of the selected accelerated corrosion test relate to every aspect of the corrosion behavior in the intended service application. Test parameters and methods must be adjusted to suit the numerous materials and corrosion protection strategies associated with exposure to gaseous, liquid, and solid environments. A nonexhaustive list of possible factors to consider in an accelerated corrosion test for any given in-service condition follows: Reactive ionic and nonionic chemical or biological species in aqueous solutions Reactive ionic and nonionic chemical or biological species in nonaqueous solutions Nonreactive but possibly participative ionic and nonionic chemical or biological species in aqueous or nonaqueous solutions Oxygen gas Ozone Solar radiation of various wavelengths Particles that abrade the material surface Particles that temporarily adhere to the material surface Particles that permanently adhere to the material surface Temperature extremes and cycles Pressure extremes and cycles Rate of replenishment of reactants to the material surface Rate of replenishment of nonreactants to the material surface Level of agitation (or stagnation) of environment Initial surface chemistry and condition of the test material Initial, intermediate, and final reaction products (soluble) Initial, intermediate, and final reaction products (insoluble, including biofilms and gels) Initial, intermediate, and final reaction products (insoluble and adherent, including biofilms and gels)

REFERENCES

Over the years, the most important insight that has come to light is that the highest rate of corrosion damage occurs during the transition from wet to dry, not when a material is fully immersed in the corrosive environment [23]. During the drying event, the concentrations of all reactants become optimum for rapid metal dissolution. Reproducing such an event in a practical laboratory accelerated corrosion test for consumer products continues to be an elusive endeavor. C3. Statistical Considerations The topic of experimental design, particularly the selection of sample size or replication, is a bitter rival to accelerated corrosion testing for the spirited scientific debate crown. As can be expected, mixing the two can result in ill-advised decisions or gridlock if not handled judiciously. Replication, randomization, and blocking are the cornerstones of all experimental designs and should receive full consideration before embarking on an accelerated corrosion testing program. Plan for the final statistical analysis long before preparing specimens for testing. Decide what will be measured (and with what instruments) so that the measuring system may be evaluated for repeatability, reproducibility, accuracy, and precision long before estimates of experimental error are needed. Decide what will be measured (and with what tools), estimate the standard deviation of the values to be measured, commit to a statistical level of significance to be met for both type I and type II errors, and determine the difference between two values that is deemed to be practically significant. With that information in hand, one can calculate the sample size (replication) necessary to achieve the stated goals. Operating characteristic (OC) curves are useful in this regard [24]. Many promising product ideas have been abandoned due to inadequate replication, resulting in the inflation of type I and type II errors during initial screening experimentation. Numerous statistical texts outline the process of experimental design and statistical data analysis [24–27].

D. CONCLUDING REMARKS An athlete may be judged in initial tryouts by his or her ability to run a given distance in a short period of time, to lift a prescribed weight numerous repetitions, or to accomplish any of a number of astonishing physical feats, but the true test of the athlete’s ability to perform comes when all of the conditions of the particular sport are present in the correct proportions and under relevant circumstances, that is, during game or match conditions. Particular physical skills, as independently measured, are mere predictors or enabling attributes that contribute to the athlete’s ability to play a given sport and may have little relevance outside of the testing environment.

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Similarly, an accelerated corrosion test or series of tests and measurements cannot, a priori, be expected to tell a complete story. Expert consideration, coaching if you will, is needed to derive the desired information from an imperfect set of raw data. Don’t just toss specimens into a test cabinet and ignore them until the test is completed—observe the degradation process by whatever means available to illuminate the characteristics that give you the information needed to make a decision. Corrosion is a process, not a property. Development, selection, conduct, and interpretation of accelerated corrosion tests benefit greatly if that fact is remembered.

REFERENCES 1. F. L. LaQue, “A Critical Look at Salt Spray Tests,” Mater. Methods, 35(2), 77–81 (1952). 2. F. L. LaQue, “A Critical Look at Salt Spray Tests,” Mater. Methods, 35(3), 77–81, 156, 158, 160, 162, 164, 166, 168 (1952). 3. W. D. McMaster, The Accelerated Corrosion Testing of Metals, First International Congress on Metallic Corrosion, Butterworths, London, 1962, pp. 679–684. 4. L. Schlossberg, “Corrosion Theory and Accelerated Testing Procedures, Part 1,” Met. Finish., 62(4), 57–63 (1964). 5. L. Schlossberg, “Corrosion Theory and Accelerated Testing Procedures, Part 2,” Met. Finish., 62(5), 93–99 (1964). 6. J. Mazia, More Metal Finishing Myths (MFM’s). Metal Finishing 75(5), 77–81 (1977). 7. T. S. Lee and K. L. Money, “Difficulties in Developing Tests to Simulate Corrosion in Marine Environments,” Mater. Perform., 23(8), 28–33 (1984). 8. W. Funke, et al., “Unsolved Problems of Corrosion Protection by Organic Coatings: A Discussion,” J. Coat. Technol., 58(741), 79–86 (1986). 9. Cleveland Society for Coatings Technology Technical Committee, “Correlation of Accelerated Exposure Testing and Exterior Exposure Sites,” J. Coat. Technol., 66(837), 49–67 (1994). 10. J. Maxted, “Short Term Testing and Real Time Exposure,” J. Corros. Sci. Eng., 2(15) (1999). 11. J. Repp, “Accelerated Corrosion Testing—Truths and Misconceptions,” Mater. Perform., 41(9), 60–63 (2002). 12. J. Guthrie, B. Battat, and C. Grethlein, “Accelerated Corrosion Testing,” AMPTIAC Q., 6(3), 11–15 (2002). 13. G. J. Jorgensen, “A Phenomenological Approach to Obtaining Correlations between Accelerated and Outdoor Exposure Test Results for Organic Material,” J. Test. Eval., 32(6), 494–499 (2004). 14. K. Barton, Acceleration of Corrosion Tests on the Basis of Kinetic Studies of the Rate Controlling Combination of Factors., First International Congress on Metallic Corrosion, Butterworths, London, 1962; pp. 685–690.

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15. D. Grossman, “More Realistic Tests for Atmospheric Corrosion,” J. Protec. Coat. Linings, 13(9), 40–45 (1996). 16. B. Boelen, B. Schmitz, J. Defourny, and F. Blekkenhorst, “A Literature Survey on the Development of Accelerated Laboratory Test Methods for Atmospheric Corrosion of Precoated Steel Products,” Corros. Sci., 34(11), 1923–1931 (1993). 17. A. Forshee, “Accelerated Corrosion Testing, Part 1: An Overview of 20 Possible Accelerated Corrosion Tests,” Met. Finish., 91(9), 51–54 (1993). 18. SAE, J2334, Laboratory Cyclic Corrosion Test, SAE International, Warrendale, PA, 2003. 19. F. Altmayer, “Choosing an Accelerated Corrosion Test,” Met. Finish., 100(1A), 572, 574–578 (2002). 20. H. D. Hilton, “Selecting a Cyclic Corrosion Test Cabinet,” Mater. Perform., 42(2), 76–79 (2003).

21. R. Baboian, (Ed.), Corrosion Tests and Standards: Application and Interpretation, Manual MNL 20, American Society for Testing and Materials, Philadelphia, 1995. 22. D. L. Jordan, “Location and Identity of the Cathodic Reaction During Underfilm Corrosion of Painted Galvanized Steel”, Dissertation, Illinois Institute of Technology, Chicago, IL, 1996. 23. M.Stratmann andH. Streckel, “On the Atmospheric Corrosionof Metals Which are Covered with Thin Electrolyte Layers—II. Experimental Results,” Corros. Sci., 30(6/7), 697–714 (1990). 24. M. G. Natrella, Experimental Statistics, Handbook 91, National Bureau of Standards, Washington, DC, 1963. 25. W. Nelson, Accelerated Testing, Wiley, New York, 1990. 26. D. C. Montgomery, Design and Analysis of Experiments, 2nd ed., Wiley, New York, 1984. 27. G. E. P. Box, J. S. Hunter, and W. G. Hunter, Statistics for Experimenters, 2nd ed., Wiley, Hoboken NJ, 2005.

74 HIGH-TEMPERATURE OXIDATION—TESTING AND EVALUATION C. A. C. SEQUEIRA Instituto Superior Tecnico, Lisboa, Portugal

A. Introduction B. Spectroscopy C. Conclusions References

A. INTRODUCTION One of the key parameters in high-temperature oxidation is the parabolic rate constant. This is true as long as protective oxidation determines the material behavior. Consequently, measurement of the weight change of the specimen is the key experimental technique in high-temperature oxidation. Generally, there are two possibilities. The first is to take a number of specimens of the same type, expose them to the respective atmosphere in a closed furnace in which defined atmospheres can be established, and take the specimens out after different oxidation states. Before and after the tests, the specimens are weighed on a high-resolution laboratory balance and the weight change Dm provided by the surface area of the specimen, A, is plotted versus time. Therefore, before exposure, the surface area A of the specimens has to be determined accurately. If the results are plotted in a parabolic manner, that is, Dm/A is squared while t remains linear, the rate constant can be determined directly from the slope, as discussed in Chapter 20. A more elegant method is to use continuous thermogravimetry [1]. In this case, a platinum or quartz string is attached to a laboratory balance that extends down into a furnace. At the lower end of the string, a coupon specimen is attached for which the surface area must be determined

before the test. A movable furnace can also be installed that allows thermocyclic oxidation testing by periodically moving the furnace over and away from the specimen area. Ideally, the specimen should be in a quartz chamber or a chamber of another material that is highly corrosion resistant so that defined gas atmospheres can be used in the tests. The interior of the microbalance must be shielded against the aggressive gas atmospheres, usually by a counterflux of a nonreactive gas, such as argon. In a more sophisticated type of thermobalance, acoustic emission measurements can be made, for example, by acoustic emission thermography (AET) [2]. This becomes possible if a waveguide wire is attached to the string hanging down from the balance, to which the specimen is attached. In particular, under thermocyclic conditions, acoustic emission measurements allow the determination of the critical conditions under which the oxide scales crack or spall [3]. If this type of scale damage is accompanied by mass loss due to spallation of the scale, then this is directly reflected in the mass change measurements and can be correlated with acoustic emission results. In some situations, internal oxidation or corrosion may occur which cannot be detected directly by thermogravimetric measurements. Therefore, it is necessary to perform metallographic investigations as well. In particular, for continuous thermogravimetric testing, at the end of each test a metallographic cross section should be prepared in order to check whether the mass change effects measured in the tests are caused by surface scales alone or whether the metal cross section has been significantly affected. Furthermore, if the kinetics of internal corrosion are to be determined, it is necessary to perform discontinuous tests where specimens are taken out of the test environment after different testing times and then investigated by metallographic techniques [1, 4–7].


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HIGH-TEMPERATURE OXIDATION—TESTING AND EVALUATION

Standard high-temperature corrosion investigations usually also include an analysis of the corrosion products formed in the tests or under practical conditions because this allows conclusions as to which are the detrimental species in the environment and whether protective scales had formed. In most cases, this is done either in the scanning electron microscope by using energy-dispersive X-ray analysis or with metallographic cross sections in the electron probe microanalyzer (wavelength-dispersive X-ray analysis). Another tool may be X-ray diffraction by the grazing angle (GAXRD) technique, which allows analysis of the composition of thin layers. Common experimental investigation techniques used to assess corrosion morphology, identify corrosion products, and evaluate mechanical properties are the following [1, 4, 8–11]: (a) Corrosion morphology assessment Optical microscopy Scanning electron microscopy (SEM) Electron probe microanalysis (EPMA) Transmission electron microscopy (TEM) (b) Corrosion product identification X-ray diffraction (XRD) Energy-dispersive X-ray analysis (EDX) Secondary ion mass spectroscopy (SIMS) X-ray photoelectron spectroscopy (XPS) Auger electron spectroscopy (AES) Laser Raman spectroscopy (LRS) (c) Mechanical properties evaluation

B. SPECTROSCOPY Chemical analysis by spectroscopy has made rapid advances in high-temperature studies and almost always includes equipment for high-resolution microscopy. Several books and monographs are available, including most of the old and newly developed techniques [10, 12–14]. Glow discharge spectroscopy (GDMS) is fast, sensitive, accurate, simple, and reliable and can be used for surface analysis if the specimen can be attached to a vacuum cell [15]. TABLE 74.1. Methods of Material Characterization by Excitation and Emission Primary Excitation

Detected Emission

Methods of Analysis: Name and Nomenclature

Photons, optical

Optical

Spectroscopy AA: Atomic absorption ) IR: Infrared ) UV: Ultraviolet ) Visible ) UPS: Vacuum UV photoelectron spectroscopy Outer shell XPS: X-ray photoelectron spectroscopy Inner shell; also called ESCA ESCA: Electron spectroscopy for chemical analysis XFS: X-ray fluorescence spectrometry XRD: X-ray diffraction EPMA: Electron probe microanalysis SEM: Scanning electron microscopy TEM: Transmission electron microscopy STEM: Scanning transmission electron microscopy SAM: Scanning Auger microanalysis AES: Auger electron spectroscopy SCANIIR: Surface composition by analysis of neutral and ion impact radiation IIXA: Ion-induced X-ray analysis ToFMS: Time-of-flight mass spectrometry SIMS: Secondary ion mass spectrometry IPM: Ion probe microanalysis ISS: Ion scattering spectrometry RBS: Rutherford backscattering spectrometry ES: Emission spectroscopy SSMS: Spark source mass spectrography

Electrons Photons, X-rays

X-rays

Electrons

Creep rupture Postexposure ductility Modulus of rupture (MOR) In general, creep rupture, hardness, and MOR have been used equally to assess the mechanical properties of corroded test pieces. When the material is difficult to grip (as is a ceramic), its strength can be measured in bending. The MOR is the maximum surface stress in a bent beam at the instant of failure [International System (SI) units, megapascals; centimeter–gram–second (cgs) units, 107 dyn/cm2]. One might expect this to be exactly the same as the strength measured in tension, but it is always larger (by a factor of about 1.3) because the volume subjected to the maximum stress is small and the probability of a large flaw lying in the highly stressed region is also small. (In tension all flaws see the maximum stress.) The MOR strictly applies only to brittle materials. For ductile materials, the MOR entry in the database is the ultimate strength [1]. Spectroscopic techniques used for chemical analysis of oxidation problems and characterization of thin layers of corrosion scales are of considerable importance, as discussed in the following paragraphs.

Electrons

X-rays Electrons

Ions

Optical

X-rays Ions ()

Radiation

Optical Ions ()

TABLE 74.2.

Summary of Various Characteristics of the Analytical Techniques

Characteristic

AES

XPS

ISS

SIMS

RBS

NRA

IXX

Sample alteration

High for alkali halogen organic Insulators

Low

Low

Low

Very low

Very low

Very low

Elemental analysis Sensitivity, variation, resolution Detection limits Chemical state Quantification

Good Good

Good Good

Good Fair

Poor Good

Fair Fair

Fair Good

Good Good

0.1% Yes With difficulty, req. standards 200 nm Atomic layer to Destructive, sputter

0.1% Yes With difficulty, req. standards 2 mm Atomic layer to Destructive, sputter

0.1% No With difficulty, req. standards 100 m Atomic layer to Destructive, sputter

10 4% or higher Yes Very difficult, req. standards 100-1 m Atomic layer to Destructive, sputter

10 3% or higher No Absolute, no standards 1 mm 10 nm Non destructive

10 2% or higher No Absolute, no standards 1 mm 10 nm Non destructive

10 2% or higher No Absolute, no standards 1 mm None Very difficult

Lateral resolution Depth resolution Depth analysis

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TABLE 74.3. Outline of Some Important Techniques to Study Metallic Surfaces Technique

Abbreviation

Information

Comments

Optical microscopy

OM

Surface topography and morphology

Transmission electron microscopy Scanning electron microscopy

TEM

Surface topography and morphology

SEM

X-ray photoelectron spectrometry Secondary ion mass spectroscopy Auger electron spectroscopy

XPS (ESCA)

Surface topography and morphology combined with X-ray spectroscopy gives “bulk” elemental analysis Chemical composition, depth profiling

Contact angle measurement

—

Inexpensive but modest resolving power and depth of field Very high resolution but requires replication; artifacts can be a serious problem Resolving power >> optical microscopy; preparation easier than TGEM and artifacts much less likely Especially useful for studying adhesion of polymers to metals Extremely high sensitivity for many elements High spatial resolution which makes the technique especially suitable for composition-depth profiling Inexpensive; rapid

SIMS AES

Elemental analysis in “monolayer range,” chemical composition and depth profiling Chemical composition, depth profiling, and lateral analysis Contamination by organic compounds

The resolving power in depth profiling is similar to AES and SIMS. A 1-kV glow discharge causes ion bombardment and surface erosion which is fed (optically) to a multichannel spectrometer for elemental analyses. Other long and complex methods of surface analysis, such as AES, SIMS, XPS, ion scattering spectroscopy (ISS), Rutherford backscattering (RBS), nuclear reaction analysis (NRA), ion-induced X-ray emission (IIXA), and ESCA, are difficult for field use. Several authors have reviewed these methods [16–28]. Tables 74.1–74.4 compare the techniques, and Figure 74.1 shows the relative sizes of areas analyzed using these techniques.

Commonly used methods are optical and scanning electron microscopy for surface studies. Transmission electron microscopy of interfaces has been explored. Selected area diffraction patterns show the orientation between different grains. In a ceramic coating, the interface between different phases can be coherent, semicoherent, or incoherent. Coherent phases are usually strained and can be studied by TEM contrast analysis. Other aspects of analytical electron microscopy analysis are discussed [29, 30]. TEM resolution is better than 1 nm and selected volumes of 3 nm diameter can be chemically analysed. Methods of preparing thin TEM transparent foils are described [16, 31–33].

TABLE 74.4. Types of Samples and Techniques Nearly Appropriate for Their Analysis Required Sample Analysis Depth profiling lower Z elements and thin films; trace or minor analysis of light elements; quantitative analysis Depth profiling of higher Z elements and thin films; trace analysis of heavy elements in light matrix; quantitative analysis Trace, minor, and major element analysis in thicker samples; quantitative analysis Minor and major elements at surface or interface of small samples Trace elements at surface or interface of medium to small size samples; analysis of insulators; sputter profiling of light elements Chemical state analysis; analyses of organics, insulators Analysis of outer atom layer; analysis of insulators

Appropriate Technique NRA

RBS

b

c

a

IIX

b

d e

AES SIMS

XPS ISS

Diameter

Technique

a:2.5 mm XPS b:1 mm RBS,NRA c:100 microns ISS & Micro-SIMS d:1 microns Micro-SIMS e:0.2 microns AES

FIGURE 74.1. Schematic illustrating relative sizes of areas scanned by spectrometric analytical techniques.

CONCLUSIONS

Photoemission with synchrotron radiation can probe surfaces on an atomic scale [34, 35], but this method requires expensive equipment. Complex impedance measurements can separate surface and bulk effects, but problems of interpretation need to be resolved [10, 34]. X-ray and gamma radiographs, as used in weld inspection, can be used to inspect coatings for defects. The method has been discussed by Helmshaw [36]. Inclusions, cracks, porosity, and sometimes lack of fusion can be detected. Surface compositions of ion-implanted metals have been studied by RBS [17, 35]. In this nondestructive way a microanalysis of the near-surface region is obtained. Interpretation is relatively easy. Assessment of radiation damage in ion-implanted metals by electron channelling is described using SEM [35]. NRA for the characterization of surface films is described [10, 35]. AES and XPS analyze the top of the surface only and erosion by ion bombardment or mechanical tapering is needed to analyze deeper regions. AES detects 0.1% of an impurity monolayer in a surface. Auger electrons are produced by bombarding the surface with low-energy (1–10 keV) electrons. In XPS the surface is exposed to a soft X-ray source and characteristic photoelectrons are omitted. Both AES and XPS electrons can escape from only 1 nm depth from the surface, and so these are surface analytical methods [10, 17]. It is most important to avoid surface contamination during preparation for surface analysis. Semiquantitative in situ analysis by AES has been reported [37]. Nitrides and other compound refractory coatings are frequently analyzed by AES and RBS methods. Depth and crater edge profiling have been done for direct-current (dc) magnetron sputtered and ARE (activated reactive evaporation) samples of (Ti, Al) N, TiN, and TiC coatings [38, 39]. Round-robin tests of characterization by including a range of analyses such as XPS, EPMA, XRS, AES, Atom Probe Microanalysis (APMA), and XRD are not uncommon. Among these, XRD was felt to be unreliable [40]. Ion spectroscopy is a useful technique for surface analysis [10, 17]. ISS uses low-energy back-scattered ions [41] and has a high sensitivity. SIMS has the possibility of sputter removal of layers allowing depth profiling [17]. It can act as a stand-alone single system to solve surface analysis. Three-dimensional SIMS of surface-modified materials and examination of ion implantation is reported [42]. Lattice vacancy estimation by positron annihilation is another approach [10, 43]. Transmission and scanning electron microscopy (TEM and SEM) are valuable techniques, and replication methods, using, for example, acetate replicas, can nondestructively reveal surface features of specimens too thick for TEM [1, 17]. ARE coatings of V–Ti in C2H2 give wear-resistant (V,Ti)C coatings. The hardness is related to grain size, stoichiometry, free graphite, and cavity networks. SEM and XRD analysis could not be used to explain the large hardness variations obtained by varying temperature and gas pressure, but TEM revealed

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microstructural changes [1, 10, 16]. Beta backscatter and X-ray fluorescence have low sensitivity (0.5 cm2/min and 1 cm2/h, respectively). Thickness and uniformity of silica coatings on steel have been determined by X-ray fluorescence measurements of Si concentrations along the surface [16, 44]. Round-robin tests for microstructure and microchemical characterization of hard coatings have included XPS, UPS, AES, EELS (electron energy loss spectroscopy), EDX, WDX (wave-dispersive X-ray analysis), RBS, SIMS, TEM, STEM, and XTEM (X-ray transmission electron microscopy) [1, 44–46]. Field emission STEM has been applied for profiling Y across a spinel– spinel grain boundary [47].

C. CONCLUSIONS In studying oxidation behavior at high temperatures, the foremost requirement is to monitor the extent and kinetics of attack. To obtain a complete mechanistic understanding, such data have to be augmented by precise details of all the processes involved, starting with the chemical reaction sequence leading to the formation of gaseous products and solid products at the reacting surface. The development and failure of protective surface scales crucially govern the resistance of most materials in aggressive environments at elevated temperatures. Knowledge is also essential on the changes throughout the exposures of the scale chemical composition, physical structure (including topography), stress state, and mechanical properties as well as on the scale failure sequence (e.g., by cracking and spallation). All these processes involved in high-temperature oxidation are dynamic. Therefore, to obtain unambiguous information, the main experimental approach in research should be based on in situ methods. These can be defined as being techniques which either measure or observe directly hightemperature oxidation processes, as they happen, in real time. Although numerous in situ methods have been developed, to date, with several notable exceptions, the most important being controlled atmosphere thermogravimetry, the deployment of these techniques often has been limited. This may be attributed largely to experimental difficulty and also to the lack of suitable equipment. Current understanding of the chemical and physical characteristics, stress state, and mechanical properties of oxidation scales largely derives from postoxidation investigations. In fact, certain detailed aspects, for example, variations in mechanical properties and microstructure through scales, can be revealed only by postoxidation studies. The two main experimental approaches, in situ oxidation and postoxidation, are not mutually exclusive, as they complement and augment each other. Nevertheless, at the current state of mechanistic knowledge of high-temperature oxidation, further understanding of many critical facets (e.g., the breakdown of

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protective oxide scales) will be revealed only by real-time experimentation. These requirements taken in conjunction with recent advances in both commercial and experimental equipment design/capabilities and in data storage/processing make it imperative that all investigators in this field be fully aware of the available in situ experimental test methods. The purpose of this chapter is to provide a very brief summary of the main techniques in current use, their main limitations, and scope for development. Information on the detailed methodology of any technique or the complete results of any specific study using any such technique should be obtainable from the references given to published papers. REFERENCES 1. H. J. Grabke and D. B. Meadowcroft (Eds.), Guidelines for Methods of Testing and Research in High Temperature Corrosion, Institute of Materials, London, 1995. 2. M. Walter, M. Sch€utze, and A. Rahmel, Oxid. Met., 39, 389 (1993). 3. M. Sch€utze, Protective Oxide Scales and Their Breakdown, The Institute of Corrosion, Wiley, Chichester, 1997. 4. N. Birks, G. H. Meier, and F. S. Pettit, Introduction to the High Temperature Oxidation of Metals, Cambridge University Press, Cambridge, 2006. 5. Y. Wouters, A. Galerie, and J. P. Pettit, Solid State Ionics, 104, 89 (1997). 6. B. Glaser, K. Rahts, M. Schorr, and M. Sch€utze, Sonderband des Praktischen Metallographie, 25, 75 (1994). 7. H. Baboian (Ed.), High-Temperature Gases (Corrosion Tests and Standards: Applications and Integration), American Society for Testing and Materials, Washington, DC, 1995. 8. S. Taniguchi, T. Maruyama, M. Yoshiba, N. Otsuka, and Y. Kawahara (Eds.), High Temperature Oxidation and Corrosion 2005, Trans. Tech. Publications, Z€urich, 2006. 9. A. Rahmel (Ed.), Aufbau von Oxidschichten auf Hochtemperaturwerkstoffen und ihre technische Bedeutung, Deutsche Gesellschaft f€ur Metallkunde, Oberursel, 1982. 10. P. Marcus and F. Mansfeld (Eds.), Analytical Methods in Corrosion Science and Engineering, CRC Press, Taylor & Francis Group, Boca Raton, FL, 2006. 11. C. A.C. Sequeira, Y. Chen, D.M.F. Santos, and X. Song, Corros. Prot. Mater., 27, 114 (2008). 12. P. A. Psaras and H. D. Langford (Eds.), Advancing Materials Research, National Academy Press, Washington, DC, 1987. 13. O. Kubaschewski and B. E. Hopkins, Oxidation of Metals and Alloys, Butterworths, London, 1967. 14. P. Kofstad, High Temperature Corrosion, Elsevier Applied Science, London, 1988. 15. R. Berneron and J. C. Charbonnier, in Proceedings 7th ICVM Conference, Tokyo, 1982, p. 592. 16. E. Lang (Ed.), Coatings for High Temperature Applications, Applied Science, New York, 1983.

17. D. M. Brewis (Ed.), Surface Analysis and Pretreatment of Plastics and Metals, Applied Science, London, 1982. 18. H. Nickel, Y. Wouters, M. Thiele, and J. Quadakkers, J. Anal. Chem., 361, 540 (1998). 19. L. Pfeil, J. Iron Steel Inst., 119, 501 (1929). 20. L. Pfeil, J. Iron Steel Inst., 183, 237 (1931). 21. J. Mougin, M. Dupeux, A. Galerie, and L. Antoni, Mater. Sci. Technol., 18, 1217 (2002). 22. J. L. Liu and J. M. Blakely, Appl. Surf. Sci., 74, 43 (1994). 23. W. D. Jennings, G. S. Chottiner, and G. M. Michal, Surf. Interf. Anal., ii, 377 (1988). 24. W. J. Quadakkers, J. Jedlinski, K. Schmidt, M. Krasovec, G. Borchardt, and H. Nickel, Appl. Surf. Sci., 47, 261 (1991). 25. H. J. Grabke, M. Steinhorst, M. Brumm, and D. Wiener, Oxid. Met., 35, 199 (1991). 26. D. Clemens, K. Bongartz, W. Speier, R. Hussey, and W. J. Quadakkers, Fresenius J. Anal. Chem., 346, 318 (1993). 27. H. Viefhaus, K. Hennesen, M. Lucas, E. M. Muller-Lorenz, and H. J. Grabke, Surf. Interf. Anal., 21, 665 (1994). 28. H. Bohm, Metalloberflache, 46, 3 (1993). 29. M. J. Thoma, Vac. Sci. Technol., A4, 2633 (1986). 30. H. Hansmann and J. Mosle, Adhesion, 26, 18 (1982). 31. B. E. Jacobson and R. E. Bunshah (Eds.), Films and Coatings for Technology, CEI Course, Stockolm, Sweden, 1981. 32. J. Doychak, J. L. Smialek, and T. E. Mitchell, Met. Trans. A, 20A, 499 (1989). 33. A. Strecker, U. Salzberger, and J. Mayer, Practical Metallogr., 30, 482 (1993). 34. J. Pask and A. Evans (Eds.), Surfaces and Interfaces in Ceramic Metal Systems, Plenum, New York, 1980. 35. V. Ashworth, W. A. Grant, and R. P. M. Procter (Eds.), Ion Implantation in Metals, Pergamon, London, 1980. 36. R. Helmshaw, Industrial Radiology, Applied Science, London, 1982. 37. A. Bosseboeuf and D. Bouchier, Surf. Sci., 162, 695 (1985). 38. H. A. Jehn, S. Hofman, and W.-D. Munz, Thin Solid Films, 153, 45 (1987). 39. N. Kaufherr, G. R. Fenske, D. E. Busch, P. Lin, C. Despandey, and R. F. Bunshah, Thin Solid Films, 153, 149 (1987). 40. A. J. Perry, C. Strandberg, W. D. Sproul, S. Hofmann, C. Erneberger, J. Nickerson, and L. Cholet, Thin Solid Films, 153, 169 (1987). 41. A. W. Czanderna (Ed.), Methods of Surface Analysis, Elsevier, London, 1975. 42. R. H. Fleming, G. P. Meeker, and R. J. Blattner, Thin Solid Films, 153, 197 (1987). 43. J. Brunner and A. J. Perry, Thin Solid Films, 153, 103 (1987). 44. M. J. Bennett, J. Vac. Sci. Technol., B12, 800 (1984). 45. J. E. Sundgren, A. Rockett, J. E. Greene, and U. Helmersson, J. Vac. Sci. Technol., A4, 2770 (1986). 46. S. Bose, High Temperature Coatings, Elsevier, Amsterdam, 2007. 47. M. J. Bennett and A. T. Tuson, Mat Sci. Eng., 92, 180 (1989).

75 TESTING FOR FLOW EFFECTS ON CORROSION K. D. EFIRD Efird Corrosion International, Inc., The Woodlands, Texas

A. Introduction B. Correlation of test data to operating facilities C. Flow effect test techniques D. Low-turbulence flow corrosion test methods References

A. INTRODUCTION Fluid velocity was long used as the primary parameter for scaleup of laboratory test results to field applications, but this concept began changing in the 1970s. Data relating the calculated hydrodynamic parameter of wall shear stress to corrosion were first published by Efird in 1977 for copper alloys in flowing seawater [1]. Corrosion science now understands that fluid flow must be expressed in terms broadly related to fluid flow parameters common to all hydrodynamic systems to allow application of laboratory test data to field operations [2–11]. These hydrodynamic parameters are calculated from empirical equations developed to characterize fluid flow. The hydrodynamic parameters employed are wall shear stress (tw) and mass transfer coefficient (k).

B. CORRELATION OF TEST DATA TO OPERATING FACILITIES The two primary considerations in the correlation of laboratory data to corrosion in operating facilities are the material

tested and the laboratory test procedure. The material tested must correspond as closely as possible to the material used in the facility in both alloy chemistry and metallurgical structure. The laboratory corrosion tests for flow-induced corrosion must be conducted in a manner that allows calculation of the mass transfer coefficient or wall shear stress [2–11]. Experimentally determined corrosion rates are then applied to the operating facilities for identical calculated hydrodynamic parameters. The basic assumptions are as follows: . . .

Calculation of the hydrodynamic parameters is valid. Calculated hydrodynamic parameters are those controlling corrosion or are intimately related to them. Scaleup of these hydrodynamic parameters to field operations with respect to corrosion is valid.

Both wall stress (tw) and mass transfer coefficient (k) meet these basic criteria. The process of correlating laboratory data to field operations is outlined in Figure 75.1 [12]. The operating flow rate allows calculation of the wall shear stress, which is then correlated through the wall shear stress to the measured laboratory corrosion rate. The existing applicable equations for both wall shear stress and mass transfer coefficient assume equilibrium conditions. As discussed in Chapters 17 and 64, both equilibrium and disturbed-flow, steady-state conditions exist in operating systems. While the existing equations and test methods are effectively used for equilibrium flow conditions in operating systems, they may not be applicable to disturbed-flow conditions. The basic methodology for relating an experimentally determined corrosion rate to allow prediction of the corrosion


1059

TESTING FOR FLOW EFFECTS ON CORROSION

LABORATORY DATA

WALL SHEAR STRESS

FIELD DATA

WALL SHEAR STRESS

1060

Field Wall Shear Stress Lab Wall Shear Stress

GAS

LIQUID

Field Flow Rate Lab Corrosion Rate

FIELD FLOW RATES

FIELD STEADY STATE WALL SHEAR STRESS

CORROSION RATES

CORROSION RATE SCALES TO WALL SHEAR STRESS

FIGURE 75.1. Process for relating laboratory data to facility operating systems using hydrodynamic conditions [12].

rate for flow-induced corrosion in an operating system is outlined in Figure 75.2 [13]. This provides a procedure to relate the results of a corrosion test directly to corrosion in an operation system where flow-induced corrosion conditions are expected or must be considered. The steps in the process areas follows: .

.

.

The steady-state mass transfer, as defined by the limiting diffusion current density, is measured as a function of wall shear stress using a reversible, well-characterized electrochemical reaction. The equation relating the wall shear stress or mass transfer coefficient to the limiting diffusion current density is defined. The experimental apparatus for the corrosion measurements is calibrated by measuring the limiting diffusion

Assumed equilibrium conditions

LABORATORY FLOW CONDITIONS

MEASURED CORROSION RATE IN LABORATORY TEST

CALCULATED SHEAR STRESS OR MASS TRANSFER COEFFICIENT FOR LABORATORY TEST

Equations assuming hydrodynamic equivalence FIELD FLOW CONDITIONS

PREDICTED CORROSION RATE FOR FIELD APPLICATION

CALCULATED SHEAR STRESS OR MASS TRANSFER COEFFICIENT FOR FIELD APPLICATION

Assumed equilibrium conditions

FIGURE 75.2. Relating laboratory data to field applications using equilibrium conditions [12].

.

.

current density for the same reaction as a function of a convenient flow parameter. Applying the equation relating the wall shear stress or mass transfer coefficient to the limiting diffusion current density allows calibration as a function of the flow parameter. The corrosion rates obtained in the calibrated test apparatus are related to the calculated disturbed-flow wall shear stress or mass transfer coefficient in the operating system, allowing estimation of the expected corrosion rates in the operating system for disturbed flow.

C. FLOW EFFECT TEST TECHNIQUES A number of techniques are available for testing the effects of fluid flow on materials. These techniques fall into two categories, rotating systems and flow systems. They are distinguished by the means used to induce flow across the test specimen. Basically, rotating systems move the test specimen in the fluid and flow systems move the fluid across the test specimen. Techniques in use are given in Table 75.1. A comparison of the various techniques for testing the effect of flow on materials is given in Table 75.2. Volume 8 of the NACE International series Corrosion Testing Made Easy is available for those interested in conducting tests for

TABLE 75.1. Materials

Techniques for Testing Effect of Flow on

Rotating Systems

Flow Systems

Rotating cylinder Rotating cage Rotating disk

Large-diameter flow loop ( 4 in.) Small-diameter flow loop ( 1.5 in.) Jet impingement

LOW-TURBULENCE FLOW CORROSION TEST METHODS

TABLE 75.2.

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Operational Comparison of Techniques for Testing Effect of Flow on Materialsa

Criteria

Small Flow Loop

Rotating Cylinder

Jet Impingement

Rotating Cage

Large Flow Loop

Fluid requirements Construction cost Operating cost Test work effort k and/or t? Multiphase testing? Scale up—liquid? Scale up—gas? Use coupons? Use ER? Use LPR, EIS?

High High High Medium Yes Yes Yes No Yes Yes Yes

Low Low Low Low Yes No No No Yes Yes Yes

Medium Medium Medium Medium Yes Yes Yes No Yes Yes Yes

Low Low Low Low No Yes No No Yes No No

Very High Very High Very High High Yes Yes Yes Yes Yes Yes Yes

a

Not including rotaing disk.

the evaluation of the effects of flow on the corrosion of materials [14]. This volume provides working details for the various test techniques as well as a comparison of strengths and weaknesses.

D. LOW-TURBULENCE FLOW CORROSION TEST METHODS There are no standardized corrosion test methods or test protocols for low-flow corrosion. However, based on knowledge of the low-flow conditions that can result in low-flow

corrosion, certain criteria for test protocol design are required. The objective is to model the conditions that exist in the system stratified flow or liquid holdup locations. In general, this can involve exposure of the test material to solid, fluid, and gas phases that could be present as well as the respective interfaces. This can be as simple as a horizontal coupon with solids placed on the coupon or as complex as a single coupon that passes through all the existing phases. This test arrangement has the benefit of including any galvanic effects that might be present. An example of this setup is shown in Figure 75.3.

FIGURE 75.3. Example of a low-flow test using a single coupon that passes through all the existing phases.

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TESTING FOR FLOW EFFECTS ON CORROSION

REFERENCES 1. K. D. Efird, “The Effect of Fluid Dynamics on the Corrosion of Copper Base Alloys in Seawater,” Corrosion, 33(1), 3–8 (Jan. 1977). 2. T. Y. Chen, A. A. Moccari, and D. D. Macdonald, “The Development of Controlled Hydrodynamic Techniques for Corrosion Testing,” Paper No. 292, CORROSION/91, National Association of Corrosion Engineers, Cincinnati, OH, Mar. 11–15, 1991. 3. E. Heitz, “Chemo-Mechanical Effects of Flow on Corrosion,” MTI Publication No. 23, MTI Project No. 15, Materials Technology Institute, Columbus, OH, 1986. 4. B. T. Ellison and C. J. Wen, “Hydrodynamic Effects on Corrosion,” in Tutorial Lectures in Electrochemical Engineering and Technology, R. Alkire and T. Beck (Eds.), AIChE Symposium Series, Vol. 77, AIChE, New York, 1981, pp. 161–169. 5. K. D. Efird et al., “Experimental Correlation of Steel Corrosion in Pipe Flow with Jet Impingement and Rotating Cylinder Laboratory Tests,” Corrosion, 49(12), 992 (Dec. 1993). 6. K. D. Efird et al., “Wall Shear Stress and Flow Accelerated Corrosion of Carbon Steel in Sweet Production,” in Proceedings: 12th International Corrosion Congress, Houston, TX, Sept. 19–24, 1993.

7. J. A. Herce et al., “Effects of Solution Chemistry and Flow on the Corrosion of Carbon Steel in Sweet Production,” Paper No. 95111, CORROSION/95, Orlando, FL, Mar. 1995. 8. K. G. Jordan and P. R. Rhodes,“Corrosion of Carbon Steel by CO2 Solutions: The Role of Fluid Flow,” Paper No. 95125, CORROSION/95, Orlando, FL, Mar. 1995. 9. C. DeWaard, U. Lotz, and A. Dugstad,“Influence of Liquid Flow Velocity on CO2 Corrosion: A Semi-Empirical Model,” Paper No. 95128, CORROSION/95, Orlando, FL, Mar. 1995. 10. B. F. M. Pots, “Mechanistic Models for the Prediction of CO2 Corrosion Rates under Multi-Phase Flow Conditions,” Paper No. 95137, CORROSION/95, Orlando, FL, Mar. 1995. 11. W. P. Jepson, S. Bhongale, and M. Gopal, “Predictive Model for Sweet Corrosion in Horizontal Multiphase Slug Flow,” Paper No. 96019, CORROSION/96, Denver, CO, Mar. 1996. 12. K. D. Efird, “The Effect of Disturbed Flow on Flow Accelerated Corrosion,” EPRI Corrosion and Degradation Conference, St. Petersburg Beach, FL, June 2–4, 1999. 13. K. D. Efird, “Disturbed Flow and Flow Accelerated Corrosion in Oil and Gas Production,” J. Energy Resources Technol., 120(1), 72–77 (Mar. 1998). 14. P. Roberge, “Erosion-Corrosion,” in Corrosion Testing Made Easy, Vol. 8, B. C. Syrett (Ed.), NACE International, Houston, TX, 2004.

76 ACCELERATED TESTING OF ELECTRONICS TO SIMULATE LONG-TERM WORLDWIDE ENVIRONMENTS L. F. GARFIAS-MESIAS DNV Columbus, Inc., Dublin, Ohio

M. REID Stokes Research Institute, University of Limerick, Limerick, Ireland

A. B. C. D.

Introduction Mixed flowing gas testing Corrosion of electronic equipment worldwide Accelerated corrosion testing to simulate worldwide corrosion of electronics E. Summary References

A. INTRODUCTION Because of the very high rate at which new commercially available electronic materials and devices have been introduced into the market and phased out in the past 50 years, it has been difficult to understand the mechanisms and factors that influence their corrosion performance. In the late 1990s, research to develop new electronic materials and devices typically took one to two years, followed by a very fast ramp-up to production and worldwide deployment that typically took another year. The equipment that contained these devices became obsolete by their fifth year in service, since new, faster and higher capacity equipment was needed to keep up with the demands for modern communications. In this very fast cycle of supply and demand of new equipment with highperformance electronic devices, failure of electronics was traditionally regarded as nonexistent, mainly because suppliers

and users were willing to exchange dated technology that showed poor performance for the latest advances in technology. As a consequence, there was no time to investigate the failure of devices and components. This hindered the improvement of the design of the new generations of electronic devices. Another important issue in the fast-changing world of electronics was globalization. In the 1960s and 1970s, most electronic equipment had to be tested to stringent requirements. Qualification of electronic devices took at least a couple of years, and companies heavily promoted the good performance and reliability of their equipment. However, by the late 1990s, the need to produce and export new electronic equipment worldwide together with increasing levels of competition (that drove prices to unimaginably low levels) pushed most companies to lower the number of tests to qualify new devices and in some cases to compromise long-term reliability. These events promoted the creation of new standardized tests that could accelerate (and mimic) the environmental challenges that the equipment would see in its short life span. Mixed flowing gas (MFG) testing has been used in North America and Europe for several years to simulate the corrosion of electronics in harsh environments. This type of laboratory testing with a controlled environment is briefly explained in Section B. A few of the challenges that the industry has found, particularlywhentryingtodeployelectronicequipmentworldwide in environments that are more aggressive than the MFG Class III test, are described in Section C. Finally, in Section D, the future of accelerated testing in this ever-changing world of electronics is discussed, along with the attempts to mimic the very aggressive environments found in developing countries.


1063

1064

ACCELERATED TESTING OF ELECTRONICS TO SIMULATE LONG-TERM WORLDWIDE ENVIRONMENTS

B. MIXED FLOWING GAS TESTING

TABLE 76.1. Exposure

In the early 1980s, with the discovery of significant printed circuit board (PCB) and component failure modes (mainly due to corrosion), a number of firms and laboratories set out to develop accelerated corrosion test methods with known acceleration factors. The aim of such efforts was to shrink years of service in the worst type of environment into days of testing in the laboratories. The main goal of this testing was to prove that the field failure modes encountered during service could be replicated during the laboratory tests. To simulate the environment under operational conditions, the PCB and its components were exposed to different gas mixtures, temperatures, and relative humidity. IBM, AT&T, and Battelle laboratories participated in this effort [1]. The result of this work was the development of a MFG test, which is primarily a laboratory test in which the temperature, relative humidity, and concentration of gaseous pollutants are carefully defined, monitored, and controlled [2]. Since then, several internationally recognized standards, including American Society for Testing and Materials (ASTM) [3] and Telcordia Technologies [4], based on the MFG testing have been created. The standards are readily available and describe in detail the hardware as well as the protocol for testing and calibration. Figure 76.1 shows a typical design of a MFG testing chamber. In order to calibrate the environmental chamber and the test procedure, a set of copper coupons is exposed inside the chamber, and the weight gain of the coupons and thickness of the corrosion products are compared with the standard. This is probably the most critical part of the MFG testing and should be undertaken prior to testing. A set of four environmental corrosion classes were defined by the original group that developed the MFG tests, so that the corrosion in the field could be compared with the

Class H2S (ppb) Cl2 (ppb) NO2 (ppb) SO2 (ppb) % RH T ( C) I II III IV

Nominal Test Conditions for Standard MFG

10 100 200

10 20 50

200 200 200

200 200 200

Corrosion Chamber

Corrosive Gases

Filter

H2S Gas Monitors

NO2 SO2

PC Controls

Cl2

Temperature/Humidity Probe

Mixing Tee

30 30 50

corrosion under the laboratory test conditions in the chambers [1, 2]. The Class I environment corresponds to a benign environment in which the corrosion film on Cu in one year will not exceed 35 nm. This class is typically encountered in a central office with very good environmental control (mainly constant temperature and constant relative humidity as well as good filtering of the outdoor air entering the building). Class II is a mildly corrosive environment that ranges from a poorly controlled central office to a rural environment in which the environmental conditions are mild enough to produce a corrosion film on Cu, in one year, between 40 and 70 nm. Class III is a harsh corrosive environment typical of an outdoor cabinet containing electronic equipment without any control of its surroundings in a populated urban environment, in which the corrosion film on Cu in one year would reach between 80 and 400 nm. Finally, Class IV is a severe corrosive environment, such as in industrial areas or cities with a polluted environment, where the corrosion film on Cu would exceed 500 nm in one year. Table 76.1 summarizes the accepted test conditions for standard MFG exposure ranked in order of severity. Evidently, if the electronic equipment would be deployed in a Class I environment, there is no need for testing as the corrosion rate is significantly low that no failures due to corrosion are expected throughout the life span of the equipment.

Exhaust

Air Process Heater

70 70 75

Humidity Controller

Air Compressor

Air inlet

FIGURE 76.1. Schematic diagram of MFG test chamber setup. (Courtesy of The Electrochemical Society [23].)

CORROSION OF ELECTRONIC EQUIPMENT WORLDWIDE

Numerous authors have studied the effect of MFG (Class III) exposure on different parts of electronic components [5], including plated Cu contacts with precious metals [6] and the effect of exposure to MFG Class III [7] on contact resistance. Other types of studies include the diffusion of gases through the encapsulants [8] and creep corrosion on lead frames [9].

C. CORROSION OF ELECTRONIC EQUIPMENT WORLDWIDE Environmental data from different parts of the world suggest that more aggressive atmospheres exist in Asia [10], Australia [11], and South America [12], compared to the environment found in developed countries such as the United States [13] and Europe [14]. Typically, in those countries with harsher environmental conditions, the local regulations are not as strict as in the developed world. Contrary to what is expected [15], some cities in the developing world, particularly in rural areas with low population and low industrial activity, may have similar environmental conditions to those in cities in the developed world. However, in developing countries, the extremely aggressive environment in cities is caused mainly by the large concentration of population and/or high density of industrial activity [16]. The most common cause of failure in electronic components and devices (cited in the literature) is the high corrosion rates of metals due to the (relatively high) concentration of corrosive gases, typically sulfur-containing and/or chloridecontaining gases. Field studies and laboratory corrosion tests have been carried out to identify not only the corrosion

products of pure metals (such as Cu [17] and Ni [18]) but also the corrosion products of relevant engineering alloys (such as stainless steels [19] and Cu–Au–Sn [20]). However, few studies have been undertaken to explore the corrosion of electronic equipment and devices deployed in harsh environments (typically exposed to Class III or harsher environments). In the majority of cases, the corrosion rate has been expressed in terms of film thickness accumulated during one year on a given metal (this value can be obtained by extrapolating from measurements made during a long period of time through the year). Although it has been found that the corrosion films on different metals may be composed of several layers of different compositions, typically it has been accepted that there is a predominant film (or corrosion layer) that thickens as the metal is exposed to the aggressive environment. The corrosion layers of copper (one of the most common metals used in electronic equipment) has been investigated during exposure to both field and laboratory environments [11–14, 21]. Clearly, the most relevant studies are those in which metals are exposed to laboratory environments that mimic field environments, so that the corrosion data obtained in the short-term laboratory test can be extrapolated to a longer term in order to estimate the service life in the field environment. Figure 76.2 shows the corrosion film thickness (measured in nanometers of Cu2S per year) formed on Cu coupons exposed indoors and outdoors in 10 selected cities in Asia. As expected, the Cu2S film formed on the coupons outdoors is typically thicker, in some cases 10 times thicker, than the film formed on coupons exposed indoors. Unfortunately, most of the useful data related to corrosion of electronic equipment and devices in real environments have not been published, mainly because companies that

2,000

Cu2S Film (Nanometers/year)

1,800

1065

Indoor Cu2S Film

Outdoor Cu2S Film

1,600 1,400 1,200 1,000 800 600 400 200 -

l l l l l l al al al al ta ta ta ta ta r ia st st ur ur as as as as as st R R o o o o o oa oa u , a d 4 l/C l/C l /C l/C l/C l/C l /C si In A A ia ia ia ia ia ia ia a SI tr tr tr tr tr tr tr si s s s s s s s A A du du du du du du du In In In In In In In a a a a a a a i i i i i i i s s s s s s s A A A A A A A

FIGURE 76.2. Corrosion film thickness (nm/year) on copper exposed in indoor and outdoor environments in 10 selected cities in Asia.

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ACCELERATED TESTING OF ELECTRONICS TO SIMULATE LONG-TERM WORLDWIDE ENVIRONMENTS 250

TSP (µg/m 3)

200

150

100

50

st al rb a

n/ C oa

st al U

SI A A

A

In du

In du s SI A

st ria l/C oa

tr ia l

l ria A

A SI

SI A

U

A

rb an /In

SI A

U

R SI A A

du st

rb an

l

st al l/C oa

ur a

A

U SA

In

du st

U

ra

SA

R ur al

0

FIGURE 76.3. Total suspended particles (mg/m3) found in the outdoor environment of two selected U.S. cities compared with several selected cities in Asia.

experience failures use the knowledge gathered during the failure investigations to improve the design of future products. This gives them a competitive advantage while keeping all the knowledge inside their R&D laboratories. The most common type of data that has been published is mainly in coupons made of the same materials that are typically used in the electronic industry. However, the corrosion rate of coupons (as discussed in the next section) is very different from a device that conducts an electrical signal and that is adjacent to a dissimilar metal (most likely fabricated under different circumstances and with different materials). Most of the data available in the literature regarding failure of electronic equipment and devices have been related to corrosion of integrated circuits, connectors, resistors, transistors, or even conducting traces within circuit boards. Few studies on electronic equipment failure reported in the literature have dealt with the effect on failure of the accumulation of hygroscopic dust on the surface in a wet environment. Figure 76.3 shows the concentration of particles in two selected cities in the United States compared to several cities in Asia. The last two cities in Asia, identified as “industrial/ coastal” and “urban/coastal,” are not far from each other. More importantly, as can be seen in the graph, the concentration of dust in those cities is extremely high. Under these conditions, most electronic equipment will be covered in dust in a few months. The presence of high levels of dust [22] and high relative humidity [11] can increase the failure rate of electronics. In the ideal case where there are no gases (such as H2S, SO2, and Cl2), the main mechanism for failure of electronics may be attributed to the adsorption of water from the environment on the surfaces covered with dust, facilitating a short circuit in the powered components or even in

adjacent pins (or leads). Although some electronic devices may be protected with a conformal coating (that can retard the diffusion of water and gases through the coating), this protection may be only temporary (depending on the coating) because the device may be subjected to long periods of high humidity due to the highly hygroscopic nature of the dust layer, absorbing the water from the environment. In reality, every electronic device is surrounded by an environment containing, to some extent, hygroscopic particulates, corrosive gases, and atmospheric water (daily wet and dry cycles). The simple effect of the gases alone (together with the conservative values of temperature and relative humidity) may not be the most viable method to mimic the environmental degradation in these very aggressive conditions. In developed countries, and due to stricter regulations, the concentration of particulates (dust) is low, and the assumption that the components are not covered in dust and will not influence the corrosion of the metals is justifiable. However, in developing countries, where the concentration of particles and gases can be relatively high (sometimes 10 times greater outdoors compared to indoors), this assumption may lead to an incorrect conclusion with respect to laboratory testing and performance in the in-service environment. Despite all the efforts to understand the corrosion rates and the formation of the different corrosion layers [23] and the effect of the environment on those materials (particularly for electronic applications [24]), very little has been done to take into consideration environmental data from developing countries to develop more realistic accelerated atmospheric corrosion tests (AACTs) that will mimic the environments typically found in developing countries. This subject will be discussed in more detail in the next section.

ACCELERATED CORROSION TESTING TO SIMULATE WORLDWIDE CORROSION OF ELECTRONICS

During the decade of the 1990s, professional organizations, including ASTM, Electronic Industries Association (EIA), International Electrotechnical Commission (IEC), and Telcordia, began to standardize these test methods and published corresponding documents as guidelines [3, 4]. The standards developed for equipment deployed in the North America Region (NAR) are considered to provide accelerated aging conditions for NAR- and European-type environments only. Typically these environments are Class I or Class II and in some extreme circumstances Class III. The MFG Class III is a reliable test to accelerate these types of environmental conditions. Typical methods to evaluate the corrosion rate of metal coupons (or electronic devices) after testing include weight loss (mass change), cathodic reduction, and microscopy of the different corrosion layers. Figure 76.4 shows a comparison of the results from Cu coupons that were exposed to MFG Class III and were further evaluated using these three different methods. Figure 76.5 shows focused ion beam (FIB) images of the corrosion films on Cu coupons that were exposed to MFG Class III and further evaluated using these three different methods. Notice that the oxide film thickens with time. Although the MFG Class III test is a reliable test, care should be taken especially when selecting the test methodology to evaluate the coupons, tested devices, or field failures.

600 500 Weight Gain (µg/cm2)

D. ACCELERATED CORROSION TESTING TO SIMULATE WORLDWIDE CORROSION OF ELECTRONICS

1067

400 300 200 Cu weight gain Calculated weight gain from cathodic reduction Calculated weight gain from FIB cross sections

100 0 0

2

4

6

8

10 12 14 16 18 20 22 24

Exposure Time (days)

FIGURE 76.4. Weight gain as a function of time for Cu samples exposed to MFG Class III environment. Comparison of the weight gain measured in the Cu coupons after exposure with the weight gain calculated using the data from cathodic reduction and data observed using cross-sectional examinations (using focused ion beam). (Courtesy of The Electrochemical Society [23].)

Figure 76.4 shows that the calculated cathodic reduction corrosion layer thickness, the FIB cross-sectional measurements, and the Cu weight gain are in good agreement within the first week of exposure (when the corrosion layer is thin and well adhered to the substrate). However, after one week of exposure (or 1-mm corrosion layer) [23], the thicknesses determined by the three methods show significant differences. The discrepancies between the measurements and the

FIGURE 76.5. Cross sections of corrosion films formed on Cu coupons after exposure to MFG Class III for 1, 5, and 10 days. (Courtesy of The Electrochemical Society [23].)

1068 TABLE 76.2.

Typical Setup of Mixed Flowing Gas (MFG) for Outdoor Environments: Level 3MFG Test

AACT in aggressive environments Telcordia Telcordia EIA 1997 IEC 1995 Battelle

Test Duration (days)

Life Prediction Equivalent (years)

Temperature ( C)

RH (%)

H2S (ppb)

Cl2 (ppb)

NO2 (ppb)

SO2 (ppb)

VER (Times/h)

40

90

2000–4000

20

200

200

4 or 5

14

20

30 30 30 30 30

70 70 75 75 75

100 100 100 100 100

20 20 20 20 20

200 200 200 200 200

200 200 200 0 0

20 20 N/A 3.2 6

10 14 N/A 10 14

15 20 N/A 10 7

Abbreviations: VER, volume exchange rate inside the chamber; EIA, Electronic Industries Alliance; IEC, International Electrotechnical Commission.

Comments Round-robin tests required!

SO2 (for test IIIA)

SUMMARY

calculated values (using cathodic reduction thickness) may be attributed to a number of factors. First, variations in corrosion product thickness caused by localized differences (and stress, particularly around the coupon edges), which are not taken into consideration during FIB cross-sectional analysis, are very important. Second, in an attempt to compare the different corrosion products and film thicknesses after different exposure times, a current density value (0.35 mA/cm2 in the case of Fig. 76.4) has to be chosen. This relatively high current density may have attributed to the film reduction proceeding at an unreasonable rate with disbonding of thick corrosion products during the cathodic reduction test. Additionally, flaking of the thick corrosion product after 10 or more days (corrosion products larger than 1 mm thick) does not allow full reduction using cathodic reduction, particularly when the disbonded film is not in electrical contact with the substrate. Finally, in some cases, the corrosion products with the lowest reduction potentials are reduced first; in the case of copper, CuO is reduced first, followed by Cu2O, CuS, and finally Cu2S [25]. As copper has a naturally thin oxide film on its surface, it is reasonable to assume that for thicker films with a multilayered structure the reduction of the inner Cu oxides will invariably be reduced first. This may reduce the stability of CuS and Cu2S layers and may cause unnecessary flaking; also this effect is more likely to be more pronounced after more prolonged exposure, where the corrosion products are thicker. In most developing countries, urban cities, where most of the electronic equipment is deployed, fall into Class III and in some cases exceed Class IV [10, 15, 16]. The environmental and corrosion data as well as the few reported failures of electronic equipment in the field in different parts of the world suggest that current AACTs do not mimic the environments typically found in developing countries. In those countries, harsher environmental conditions have caused high corrosion rates of metals, leading to failure of devices and electronic equipment. In some regions of the world, corrosivity has been described in a corrosivity map [26], which indicates, by color, the typical corrosion rates, from negligible to mild and severe, measured in different regions indicated on the map. The levels of H2S (one of the most common gases found in very aggressive environments in developing countries) in the MFG for the outdoor tests (Class II or Class IV) are considerably below those levels typically found in some Asian countries [10, 15, 16]. In a recent study, Wattanabe and co-workers [10] measured the concentration of H2S near the Noboribetsu hot springs area of Hokkaido, Japan, and found that the H2S concentration ranged from 160 to 1770 ppb. Similarly, it has been reported that in parts of China, Taiwan, and Mexico the levels of H2S were as high as 850 ppb [15]. These findings strongly suggest the importance of increasing the H2S concentration in order to use MFG testing to qualify electronic equipment for use in aggressive environments. Table 76.2 shows the typical concentrations

1069

used for most mixed flowing gases according to the different standards. Based on the findings of recent studies, it is suggested that the levels of H2S when using the MFG test to qualify electronic equipment that will be deployed in very aggressive outdoor environments should be increased taking into consideration the real environment (see Table 76.2). It is clear that the suggested temperature for the MFG test (30 C) may be adequate in most cases, except when the temperature of the city (or cities) where the equipment is going to be deployed will be significantly different. In tropical cities, where the temperature year round is high, the 30 C temperature for the MFG tests seems reasonable. However, in colder areas, where there is the possibility that the temperature fluctuates significantly during the day (hot) and the night (when condensation may occur), it is recommended that these cycles be considered in the design of the test. Because the relative humidity in aggressive environments around the world ranges from low values in the northern hemisphere to high values in the tropics, the relative humidity used in the test depends on whether the objective of the test is to mimic indoor or outdoor environments. Because the outdoor environment is more humid near tropical areas in the coastal cities, the relative humidity inside the chamber should be increased; values around 90% may be acceptable if the objective is to mimic the outdoor environment. Other variables for the test should also be considered, for example, different gases, the volume exchange ratio in the chamber, and test duration sufficient for lifetime prediction. Although these variables are important, they may need to be adjusted after the values of the key variables (H2S content, temperature, and relative humidity) have been fixed. One additional variable that has not been considered during most of the AACT studies is the voltage across the electronic devices. Typically, coupons and electronic devices have been tested without power. The devices are subjected not only to particulate, gases, temperature cycles, and high relative humidity but also to a voltage. A common practice in the industry is to use coupons to measure the leakage current across the traces (lines) while the coupon is subjected to a voltage. The two voltages usually selected are 24 and 48 V because these voltages are widely used in the electronics industry.

E. SUMMARY A review of the currently available literature and common industry practices suggests that MFG testing for both indoor (typically Class II) and outdoor (normally associated with Class III) environments is the most common testing used to qualify materials and devices for use in aggressive environments. In order to mimic the overall environment in more aggressive areas of the world (e.g., in developing countries or near industrial sites in developed countries), addition of

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ACCELERATED TESTING OF ELECTRONICS TO SIMULATE LONG-TERM WORLDWIDE ENVIRONMENTS

higher levels of pollutants (including atmospheric particles and gases) should be considered for tests that are used to qualify electronic materials and devices. “Dusting” followed by AACT that includes electrical testing of interdigitated coupons may also be used as a qualification method. This may have a considerable impact when trying to reproduce field failures (e.g., creep corrosion). At the same time, it is extremely important to reproduce (and induce) during testing the same failure modes encountered in electronic devices. A comparison of available data from the United States and developing countries shows that different testing criteria should be developed (if feasible) during qualification testing for each particular environment. The ultimate goal should be to develop a meaningful AACT that can mimic the real, in-service environment while achieving an adequate acceleration factor. The available data and field experience suggest that more aggressive testing may be desirable in some situations.

REFERENCES 1. W. H. Abbott, “The Development and Performance Characteristics of Mixed Flowing Gas Environment,” IEEE Trans. Components Hybrids Manufact. Technol., 11, 22–35 (1988). 2. W. H. Abbott, “The Corrosion of Copper and Porous Gold in Flowing Mixed Gas Environments,” IEEE Trans. Components Hybrids Manufact. Technol., 13(1), 40–45 (1990). 3. ASTM B 845-97. Standard Guide for Mixed Flowing Gas Tests for Electrical Contacts, ASTM International, West Conshohocken, Pennsylvania, 2003. 4. NEBS Requirements: Physical Protection, Generic Requirements GR-63-CORE, Method 5.5.2, Telcordia Technologies, Chester, 2002. 5. D. C. Abbott, “Nickel Palladium Finish for Leadframes,” IEEE Trans. Components Packaging Technol., 22, 99–103 (1999). 6. R. J. Geckle and R. S. Mroczkowski, “Corrosion of Precious Metal Plated Copper-Alloys Due to Mixed Flowing Gas Exposure,” IEEE Trans. Components Hybrids Manufact. Technol., 14, 162–169 (1991). 7. R. Martens and M. G. Pecht, “An Investigation of the Electrical Contact Resistance of Corroded Pore Sites on Gold Plated Surfaces,” IEEE Trans. Adv. Packaging, 23, 561–567 (2000). 8. C. Hillman, B. Castillo, and M. Pecht, “Diffusion and Absorption of Corrosive Gases in Electronic Encapsulants,” Microelectron. Reliabil., 43, 635–643 (2003). 9. P. Zhao and M. Pecht, “Mixed Flowing Gas Studies of Creep Corrosion on Plastic Encapsulated Microcircuit Packages with Noble Metal Pre-plated Leadframes,” IEEE Trans. Device Mater. Reliabil., 5, 268–276 (2005). 10. M. Watanabe, H. Hirota, T. Handa, N. Kuwaki, and J. Sakai, “Atmospheric Corrosion of Cu in an Indoor Environment with a High H2S Concentration,” 17th International Corrosion Congress: Corrosion Control in the Service of Society, Las Vegas, NV, Oct. 6–10, 2008.

11. I. S. Cole, W. D. Ganther, J. D. Sinclair, D. Lau, and D. A. Paterson, “A Study of the Wetting of Metal Surfaces in Order to Understand the Processes Controlling Atmospheric Corrosion,” J. Electrochem. Soc., 151, B627–B635 (2004). 12. J. R. Vilche, F. E. Varela, E. N. Codaro, B. M. Rosales, G. Moriena, and A. Fernandez, “A Survey of Argentinean Atmospheric Corrosion: II—Cu Samples”, Corros. Sci., 39(4), 655–679 (1997). 13. D. W. Rice, P. Peterson, E. B. Rigby, P. Phipps, R. J. Cappell, and R. Tremoureux, “Atmospheric Corrosion of Cu and Silver”, J. Electrochem. Soc., 128(2), 275–284 (1981). 14. J. Tidblad and C. Leygraf, “Atmospheric Corrosion Effects of SO2 and NO2, a Comparison of Laboratory and Field-Exposed Copper,” J. Electrochem. Soc., 142(3), 749–756 (1995). 15. L. F. Garfias-Mesias, J. P. Franey, R. P. Frankenthal, and W. D. Reents, Gordon Research Conference on Corrosion, New London, NH, July 26, 2004. 16. L. F. Garfias-Mesias, J. P. Franey, R. P. Frankenthal, R. Coyle, and W. D. Reents, in CORROSION 2005, Research In Progress, NACE, Houston, TX, Apr. 4–6 2005. 17. M. Lenglet, J. Lopitaux, C. Leygraf, I. Odnevai, M. Carballeira, J.-C. Noualhaguet, J. Guinement, J. Gautier, and J. Baissel, “Analysis of Corrosion Products Formed on Copper in CI2/H2S/NO2 Exposure,” J. Electrochem. Soc., 142(11), 3690–3696 (1995). 18. S. Zakipour, J. Tidblad, and C. Leygraf, “Atmospheric Corrosion Effects of SO2, NO2, and O3, A Comparison of Laboratory and Field Exposed Nickel” J. Electrochem. Soc., 144(10), 3513–3517 (1997). 19. G. Herting, I. Odnevall Wallinder, and C. Leygraf, “A Comparison of Release Rates of Cr, Ni, and Fe from Stainless Steel Alloys and the Pure Metals Exposed to Simulated Rain Events,” J. Electrochem. Soc., 152(1), B23–B29 (2005). 20. S. Zakipour and C. Leygraf, “Evaluation of Laboratory Tests to Simulate Indoor Corrosion of Electrical Contact Materials,” J. Electrochem. Soc., 133, 21–30 (1986). 21. M. Reid, J. Punch, C. Ryan, L. F. Garfias, S. Belochapkine, J. P. Franey, G. E. Derkits, and W. D. Reents, J. Electrochem. Soc., 154, C209 (2007). 22. R. B. Comizzoli, C. A. Jankoski, G. A. Peins, L. A. Psota-Kelty, D. J. Siconolfi, J. D. Sinclair, W. Chengen, and M. Gao, in Corrosion and Corrosion Protection, J. D. Sinclair, R. P. Frankenthal, E. Kalman, and W. Plieth (Eds.), PV 2001-22, Proceedings Series, The Electrochemical Society, Pennington, NJ, 2001, pp. 691–705. 23. M. Reid, J. Punch, L. F. Garfias, K. Shannon, S. Belochapkine, and D. A. Tanner, “Study of Mixed Flowing Gas Exposure of Cu,” J. Electrochem. Soc., 155(4), C147–C153 (2008). 24. M. Reid, J. Punch, L. F. Garfias, G. K. Grace, and S. Belochapkine, “Corrosion Resistance of Cu-Coated Contacts,” J. Electrochem. Soc., 153(12), B513–B517 (2006). 25. S. J. Krumbien, B. Newell, and V. Pascucci, J. Test. Eval., 11, 357 (1989). 26. W. Hou and C. Liang, “Eight-Year Atmospheric Corrosion Exposure of Steels in China,” Corrosion, 55(1), 66 (1999).

77 TESTING FOR ENVIRONMENTALLY ASSISTED CRACKING R. D. KANE i Corrosion LLC, Houston, Texas

A. Introduction B. Background B1. Stress corrosion cracking B2. Hydrogen embrittlement B3. Liquid-metal embrittlement C. Constant-load/deflection techniques C1. Tension tests C2. Other constant-deflection specimens D. Dynamic tests E. Fracture mechanics tests F. Definition of laboratory test environments for EAC evaluation G. Summary References

A. INTRODUCTION The study of environmentally assisted cracking (EAC) involves the consideration and evaluation of the inherent compatibility between a material and the environment under conditions of either applied or residual stress. However, this is a very broad, encompassing topic with many possible combinations of materials and environments in which EAC has been investigated and documented. EAC is also a critical problem because equipment, components, and structures are intended to be used while exposed to various environments and conditions of stress and must resist EAC over prolonged periods of service. Furthermore, the materials used in construction typically have a multitude of manufacturing and process variables that may affect their metallurgical condition and structure which, in turn,

influence resistance to EAC. Testing for resistance to EAC is one of the most effective ways to determine the interrelationships among material, environmental, and mechanical variables on the process of EAC. The proportions of this subject immediately limit attempts to make simplistic use of only a single method of testing for all cases. Factors such as material type, process history, product form, active cracking mechanism(s), loading configuration and geometry, and service environment all can have a major impact on the type of specimen and test condition to be utilized for the evaluation of EAC. The prudent approach to selection of testing methods is usually to start with a survey of previous experiences from prior investigations conducted on similar classes of materials and types of environments found by surveying the published literature. Additionally, an extensive amount of information on standardized stress corrosion cracking (SCC) testing methods, evaluation procedures, and experimental techniques is available in International Organization for Standardization (ISO) 7539 Parts 1–8 [1], American Society for Testing and Materials (ASTM) and NACE standards and existing reviews published in the literature [2]. B. BACKGROUND To better establish a basis for understanding the role that each of the various cracking processes plays in the selection of EAC test methods on a particular material, the applicable EAC processes must first be identified for the specific material and environment under consideration. Mechanisms of EAC have been debated for decades and there are still controversies over the use of specific terms and definitions. Therefore, in this chapter the discussion of the subject will


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TESTING FOR ENVIRONMENTALLY ASSISTED CRACKING

be limited to basic phenomenological descriptions of the various cracking processes with an attempt to discuss them in simple but description terms and in relation to important testing-related variables. B1. Stress Corrosion Cracking Stress corrosion cracking is the formation of embrittlement whereby cracks form in a normally sound, ductile material through the simultaneous action of a tensile stress and a corrosive environment. In most cases, SCC has been associated with the process of active path corrosion (APC) whereby the corrosive attack or anodic dissolution initiates at specific, localized sites and is focused along specific paths within the material. Crack initiation often occurs at sites of local anodic attack (e.g., pits or local crevice corrosion). In some cases, crack propagation is along grain boundaries (i.e., intergranular SCC, or IGSCC); in other cases, the path is along specific crystallographic planes within the grains (i.e., Transgranular SCC). Quite often, SCC is strongly affected by alloy composition, the bulk concentration of specific corrodent species or the local concentration on the metal surface, and usually, to a lesser degree, the stress intensity. In some cases, this latter point may make use of test methods based on fracture mechanics concepts difficult to utilize effectively due to multiple crack initiations, excessive crack branching, and tendencies for nonplanar propagation of cracks. Furthermore, corrosion film characteristics (i.e., passivation) and local anodic attack (i.e., depassivation) often serve as controlling factors in SCC crack initiation and growth. Therefore, localized corrosion can promote SCC, making exposure geometry and specimen design important factors. In many cases, mechanical straining or electrochemical inducements such as crevices or controlled potential are utilized to overcome the problems and uncertainties of SCC initiation so that the inherent resistance of the material to SCC can be obtained at reasonable test duration (see Table 77.1) [3]. B2. Hydrogen Embrittlement Hydrogen is often a byproduct of corrosion and electrochemical processes and may also be a major constituent in various service environments. During electrochemical reactions in aqueous environments, it is common for a hydrogen ion (H þ ) to combine with an electron (e ) to form atomic hydrogen on the surface of the material. The effects of hydrogen on cracking are contrasted to those of local metallic dissolution in Figure 77.1 [4]. Depending on the solution and interfacial characteristics, the hydrogen atoms formed by the corrosion process may recombine to form molecular hydrogen that can simply bubble off of the specimen surface without any further complications. However, under certain circumstances, when hydrogen

TABLE 77.1. Applied Potential for SCC in Steel Exposed to Various Service Environments Environment Nitrate Liquid ammonia Carbonate Hydroxide

Potential Range (mV SCE) 250 to 1200 400 to > þ 1500 650 to 550 1100 to 850 and þ 350 to þ 500

recombination poisons (e.g., S, P, As, Sn) are present in the environment, hydrogen recombination is retarded, promoting the absorption of atomic hydrogen into the material. Once inside the material, hydrogen can affect the mechanical performance of materials in several ways: 1. The formation of internal hydrogen blisters or blisterlike cracks at internal laminations or at sites of nonmetallic inclusions in low-strength materials. These internal cracks may propagate by a process called hydrogen-induced cracking (HIC) or hydrogen blistering. No external stress on test specimens is usually required to examine this type of cracking behavior. In some cases, however, these blister cracks may take on an alignment caused by the presence of residual or applied tensile stresses, which is referred to as stressoriented hydrogen-induced cracking (SOHIC). 2. The process of hydrogen-assisted microvoid coalescence can occur during plastic straining. This can reduce the ductility of normally ductile engineering materials while not inducing brittle cracking. 3. An extreme case of ductility loss from hydrogen is the brittle fracture of susceptible materials under applied or residual tensile stresses. This form of cracking typically results in either transgranular or intergranular cracks, depending on the material type and condition, yield strength, and processing variables and is normally referred to as HEC. With respect to HEC, most susceptible materials show a major effect of stress concentration (i.e., notches) and level of stress intensity and tend to produce failures in a relatively short time (i.e., 2300 ppm) 100 mL/min-L of solution 5.1–5.4 Synthetic seawater (ASTM D1141-52 stock solution No. 1 or 2) Not stressed

where t is specimen thickness, W is specimen width, a is crack length, and b is crack thickness. As an alternative to metallographic examination, an ultrasonic C scan can be used along with a quantitative image analysis system to quantify HIC [5, 6]. The wet-fluorescent magnetic-particle inspection (WFMPI) technique can be used to detect cracks in welds and adjacent heat-affected-zones (HAZs) on the inside surfaces of pipes or vessels. This technique was found to be more sensitive in locating HIC cracks than other inspection techniques, such as radiography or dry magnetic-particle examination. B1. Sulfide Stress Cracking The standard test conditions for the NACE TMOl77 [1] test are listed in Table 78.2. This test is frequently carried out by immersing specimens in an aqueous solution containing 5% sodium chloride (NaCl), 0.5% acetic acid (CH3COOH), saturated with H2S gas at ambient temperature and pressure. Round tensile specimens are commonly used, but, as described in Chapter 77, other types of specimens are also used. Specimens are loaded to various stress levels using calibrated proof rings or dead-weight testers. Time to failure is monitored for the 720-h test duration; “no failure” is recorded if failure does not occur within this period. After characterizing the tensile properties of a material, SSC testing is carried out under constant load at various load levels, and the times to failure are plotted in terms of applied stress versus time to failure (see Fig. 77.4 in Chapter 77). In general, an SSC threshold stress can be determined and is used as a measure of cracking susceptibility. Susceptibility is

HYDROGEN DIFFUSIVITY MEASUREMENTS

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FIGURE 78.2. HIC samples from pipe material.

strongly influenced by steel strength and hardness, as shown in Figure 78.5. A sharp increase in cracking susceptibility with increased hardness is often observed. Heat treating steels to hardness levels below that at which increased cracking susceptibility occurs has been effective in minimizing field failures in sour service. B2. Determination of Dissolved Hydrogen Concentration Causing HIC of Linepipe Steels After the HIC immersion test, some investigators immerse test coupons in glycerin or mercury-filled collectors held at 45 C. The “diffusible hydrogen at 45 C” which is evolved is collected and usually reported as milliliters of H2 (NTP)/ 100 g of steel. A number of researchers have used measurement of diffusible hydrogen as a means of assessing HIC susceptibility [7–10] and have correlated diffusible hydrogen measurements with CLR and CTR [11]. The amount of

hydrogen absorbed by a test coupon at pHth, the pH just low enough to cause HIC, is the value of Cth. Ikeda et al. [10] and Hoey et al. [8] have shown that there is a critical concentration of dissolved hydrogen in a steel, termed Cth, that must be attained in order for HIC to occur.

C. HYDROGEN DIFFUSIVITY MEASUREMENTS A well-established approach to measuring hydrogen diffusion in steels involves the use of a dual electrochemical cell with the steel specimen between the charging and oxidation cells, as shown schematically in Figure 78.6. Hydrogen atoms are generated on the side of the specimen exposed to the environment of interest in the charging cell, diffuse through the steel, and are oxidized electrochemically on the opposite side of the specimen. This side is exposed to a

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TEST METHODS FOR WET H2S CRACKING

TABLE 78.2. Solutiona

Standard Test Conditions for NACE TM0177

Test period Temperature H2S concentration pH Test solution Loading stress a

720 h (30 days) 25 3 C Saturated (>2300 ppm) Start pH 2.7, end pH 4.5 5% NaCI þ 0.5% acetic acid (CH3COOH) saturated with H2S Up to the yield strength

See [1].

FIGURE 78.5. Sour gas cracking threshold stress variation with hardness for steel. Maximum hardness can be selected as that just prior to the increase in cracking susceptibility.

FIGURE 78.3. Geometry of specimens cut from plate and from pipe.

sodium hydroxide solution and is maintained at a constant potential sufficiently anodic to oxidize the hydrogen atoms that diffuse through the steel [12]. The current in the potentiostatic circuit on the oxidizing side is a direct measure of the instantaneous rate of hydrogen permeation and, by continuous recording of that current, the effective diffusion coefficient of atomic hydrogen can be determined as well as the extent of hydrogen trapping in the steel. The HIC develops when hydrogen concentration, C0. in the steel matrix exceeds the threshold hydrogen

FIGURE 78.4. The HIC susceptibility parameters: cross section perpendicular to rolling direction.

HYDROGEN DIFFUSIVITY MEASUREMENTS

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FIGURE 78.6. Schematic of the experimental apparatus for studying hydrogen permeation.

concentration, Cth [13]. The C0 depends on alloy composition, H2S partial pressure, and pH. The Cth depends on inclusions and segregation in the matrix [10]. The distribution of hydrogen atoms is derived from the solution of Fick’s second law. In an operating pipeline, hydrogen atoms entering the steel at the internal surface diffuse through the wall and exit at the external surface, where they form hydrogen gas molecules. The driving force for this flux of diffusing hydrogen atoms is the concentration gradient between the internal surface [10, 14, 15], where CH ¼ C0H , and the external surface, where

CH ¼ 0, as shown schematically in Figure 78.7. The concentration of hydrogen atoms in an operating pipeline is assumed to decrease linearly with distance through the pipe wall. At midwall, C H ¼ 0:5C0H . A schematic graph of permeation flux versus time is shown in Figure 78.8. In order to calculate the hydrogen concentration (C0H ) on the inside wall of the pipe, the hydrogen diffusion coefficient, D, must first be calculated, either using the half-rise time (t1/2) formula or using the breakthrough time (tb) formula given by Devanathan and Stachurski [12].

FIGURE 78.7. Hydrogen atom concentration gradient in steel wall.

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TEST METHODS FOR WET H2S CRACKING

From Eq. (78.3), hydrogen concentration C0H is proportional to Imax and inversely proportional to D: thus, C0H decreases as Imax becomes smaller and as D becomes greater.

REFERENCES

FIGURE 78.8. Schematic permeation transient curve for hydrogen permeation test.

The diffusion coefficient D using the half-rise time (t1/2) formula is calculated as follows: Dðcm2 =sÞ ¼

where

L2 7:2t1=2

ð78:1Þ

L ¼ thickness(cm) of steel where permeation probe was attached t1/2 ¼ time (s) needed to obtain one-half of the steadystate current, J1 (see Fig. 78.8)

The diffusion coefficient can also be calculated using the breakthrough time (tb) expression (see Fig. 78.8): D¼

L2 15:3tb

ð78:2Þ

There are other approaches to calculating values of D [16, 17], but Eq. (78.1) is often used, and Eq. (78.2) usually results in similar values. The concentration of hydrogen on the inside steel surface, C0H , is calculated from the maximum permeation current density using the relationship Peak C0H ðmmol=cm3 Þ ¼ where

Imax L DF

ð78:3Þ

C0H ¼ concentration of hydrogen at inside pipe surface (mmol/cm3) Imax ¼ peak of current density (mA/cm2) F ¼ faraday constant (96,487 C/mol) L ¼ thickness of sample (cm) D ¼ diffusion coefficient of hydrogen in steel (cm2/s) calculated from Eqs. (78.1) and (78.2)

1. NACE Standard TMO177-96, “Laboratory Testing of Metals for Resistance to Specific Forms of Environmental Cracking in H2S Enviroments,” NACE International, Houston, TX, 1996. 2. NACE Standard TM0284-96, “Evaluation of Pipeline and Pressure Vessel Steels for Resistance to Hydrogen-Induced Cracking,” NACE International, Houston, TX, 1996. 3. ASTM G 39-99, Standard Practice for Preparation and Use of Bent-Beam Stress-Corrosion Test Specimens, ASTM Volume 03-02, ASTM, West Conshohoken, PA, 1999. 4. R. S. Treseder and T. M. Swanson, Corrosion, 24(2), 27 (1968). 5. M. Elboujdaini, M. T. Shebata, and R. W. Revie, “Performance of Pipeline Steels in Sour Service,” in Proceedings, Materials for Resource Recovery and Transport, L. Collins (Ed.), The Metallurgical Society of ClM, Calgary, August 1998, CIM, Montreal, Canada, 1998, pp. 109–127. 6. A. Ikeda, A. Nakamura, and J. Kushida, “Evaluation Method of Hydrogen-Induced Crackin Susceptibilities of Steels for Pipe Lines and Pressure Vessels,” in Proc. Int. Conf. Interaction of Steels with Hydrogen in Petroleum Industry Pressure Vessel Service, Paris, France, 1989, Martin Prager (Ed.), Publ. Materials Properties Council, New York, 1993, pp. 289–304. 7. R. W. Revie, V. S. Sastri, M. Elboujdaini, R. R. Ramsingh, and Y. Lafreniere, Report MTL 91-60(TR), CANMET, Ottawa, Canada, 1991. 8. G. R. Hoey, R. W. Revie, R. R. Ramsingh, D. K. Mak, and M. T. Shehata, Report MTL 88-41(TR), CANMET, Ottawa, Canada, 1988. 9. M. Elboujdaini, M. T. Shehata, V. S. Sastri, R. W. Revie, and R. R. Ramsingh, “Hydrogen-Induced Cracking and Effect of Non-metallic Inclusions in Linepipe Steels,” Paper No. 748, CORROSION/98, NACE International, Houston, TX, 1998. 10. A. Ikeda, T. Kaneko, T. Hashimoto, M. Takeyama, Y. Sumitomo, and T. Yamura, “Development of Hydrogen Induced Cracking Resistant Steels and HIC Test Methods for Hydrogen Sulfide Service,” in Procedings, Symposium on the Effects of Hydrogen Sulfide on Steel, Paper No.6, 22nd Annual Conference of Metallurgists, Edmonton, Alberta, Canada, S. A. Bradford (Ed.), The Metallurgical Society of CIM, Montreal, Canada, 1983. 11. G. J. Biefer and M. J. Fichera, Report 84-13(TR), CANMET, Ottawa, Canada, 1984. 12. M. A. V. Devanathan and Z. Stachurski, Proc. R. Soc., 290A, 220 (1965).

REFERENCES

13. A. Ikeda, Y. Morita, F. Terasaki, and M. Takeyama, Proceedings, 2nd International Congress on Hydrogen in Metals— Paris, 4A, 7 (1977). 14. K. Van Gelder, M. J. J. Simon Thomas, and C. J. Kroese, Corrosion, 42(1), 36 (1986). 15. M. G. Hay, “An Electrochemical Device for Monitoring Hydrogen Diffusing through Steel.” CIM Symposium on “The

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Effects of Hydrogen Sulfide on Steel,” Paper 11, Edmonton, Canada, Aug. 1983, CIM, Montreal, Canada. 16. D. Noel, C. P. Vijayan, and J. J. Hechler, Surface Coatings Technol., 28, 225 (1968). 17. J. Brogan, I. M. Austen, and E. F. Walker, “Method of Calculating the Hydrogen Diffusion Coefficient in Steel from Hydrogen Permeation Data,” British Steel Corporation, London, 1987.

79 ATMOSPHERIC CORROSION TESTING D. L. JORDAN Ford Motor Company, Dearborn, Michigan

A. Introduction B. Purpose of testing C. Types of atmospheres D. Statistical considerations E. Hardware, maintenance, and practical considerations References

A. INTRODUCTION Spectacular corrosion failures occur frequently as a result of exposure of engineering materials to aggressive aqueous and nonaqueous process fluids. Many materials of construction are not exposed to such conditions but, rather, are used in the naturally occurring atmosphere and are subject to degradation processes that limit their ability to function in the desired manner. While often plain in appearance, corrosion in the atmosphere is recognized as the single most severe form of corrosion on a tonnage basis. Knowledge of the expected corrosion performance of a material in the atmosphere is needed when designing engineered components and assemblies. The purpose of this chapter is to provide information to allow design and materials selection engineers to decide if atmospheric exposure tests are a necessary part of their function and, if so, to provide them with guidelines to conduct a successful test or to interpret existing test data.

B. PURPOSE OF TESTING Before committing to the effort and expense of long-term corrosion testing, it is important to understand and record the reasons why test data are needed. The specific application, the material attributes important to that application, and the desired form of the final data must all be considered. Atmospheric corrosion test data typically are required to determine or to predict the following types of information: Material lifetime in a specific atmosphere: How long will it last before failing to satisfy a key performance indicator? Pick the winner of a group of materials in a specific atmosphere: Is material A better than material B? Corrosion rate in a specific atmosphere: What is the rate of material thickness loss? Atmospheric corrosivity comparisons: Is atmosphere A more aggressive than atmosphere B? Examples of material applications that may require the collection or interpretation of atmospheric corrosion data are as follows: Bridge work Highway guardrails


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ATMOSPHERIC CORROSION TESTING

Utility poles Power lines Electrical switchgear cabinetry Automatic teller machine cabinetry Building siding and roofs Automobile bodies Airplane skins Refrigerator wrappers or back plates Outboard motor housings Oil derricks Mail boxes Chain link fences Stadium seats and decks The material attributes or performance indicators that may be measured to provide information for proper materials selection and life prediction include: Weight loss Thickness loss Loss of impact strength Loss of tensile strength Loss of ductility Pitting Perforation Coating life Discoloration or other measure of appearance Contact resistance The data should be reported in a format that allows it to serve the desired function of the test; comparison with existing or future data is frequently, but not always, required.

C. TYPES OF ATMOSPHERES The wide variety of possible climatic conditions, combined with a multiplicity of geometric and spatial arrangements of materials, provides an unlimited assortment of atmospheres to which materials of construction may be exposed. A given material performance in one atmosphere or application should in no way be interpreted to mean that said performance would be repeated for other conditions. Many examples that demonstrate this principle exist and provide the impetus and justification for continued testing of existing and new materials in known and untested atmospheres and applications. Despite the poorly defined target, standards organizations have used different procedures to define general classifications of various atmosphere types [1, 2]. General atmospheric

classifications are needed in order to limit the number of candidate materials for an application. Classifications provide qualitative or quantitative descriptions of the key characteristics of the atmosphere that contribute to corrosion behavior: Time of wetness Sun load [exposure to ultraviolet (UV-B) radiation] Seasonal temperature and humidity cycles Natural contaminants (sea salt, particulates) Man-made contaminants (SOx, NOx, road deicing salt, tropospheric ozone, particulates) Other characteristics that may be unique to specific locations, geometries, or applications are as follows: Proximity to contaminant point sources Shelter from direct rainfall or sun load The simplest atmospheric classification scheme is based on the primary aggressors for most metallic materials in atmospheric corrosion: moisture, fallout of industrial pollution, and the chloride ion. Accordingly, the simplest classifications are as follows: Rural (very little fallout) Industrial (primarily sulfur-containing fallout) Marine (chloride-containing fallout) A refinement of the method includes some of the location characteristics cited earlier but provides better discrimination [3]: Dampness

Temperature

Contaminants

Dry Humid Marine

Tropical Temperate Arctic

Rural Urban Industrial

One entry from each of the three columns provides a reasonable qualitative description of most common atmospheres. The international program ISO CORRAG (International Organization for Standardization Technical Committee 156, Working Group 4) comprehensively characterized and classified dozens of atmospheric locations globally by two semiquantitative methods. One concentrates on the predicted corrosivity based on time of wetness and contaminant measurements [2, 4] while the other relies upon actual corrosion

STATISTICAL CONSIDERATIONS

rate measurements on standard metal specimens [2, 5, 6]. The following descriptions are now standard: Category

Time of Wetness (%) 60

t1 t2 t3 t4 t5

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It is critical to remember that the most important material and atmospheric data are not those reported by ISO CORRAG, ASTM, or any other source, but those data that are relevant to the application of interest. For instance, it has been shown that the ISO atmospheric corrosivity categories do not pertain to many sites throughout the world, including very cold areas and those where dew formation may occur in a relative humidity of

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