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Failure Analysis of Heat Treated Steel Components
L.C.F. Canale R.A. Mesquita G.E. Totten
ASM InternationalÕ Materials Park, Ohio 44073-0002 www.asminternational.org
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Copyright # 2008 by ASM InternationalÕ All rights reserved No part of this book may be reproduced, stored in a retrieval system, or transmitted, in any form or by any means, electronic, mechanical, photocopying, recording, or otherwise, without the written permission of the copyright owner. First printing, September 2008 Great care is taken in the compilation and production of this book, but it should be made clear that NO WARRANTIES, EXPRESS OR IMPLIED, INCLUDING, WITHOUT LIMITATION, WARRANTIES OF MERCHANTABILITY OR FITNESS FOR A PARTICULAR PURPOSE, ARE GIVEN IN CONNECTION WITH THIS PUBLICATION. Although this information is believed to be accurate by ASM, ASM cannot guarantee that favorable results will be obtained from the use of this publication alone. This publication is intended for use by persons having technical skill, at their sole discretion and risk. Since the conditions of product or material use are outside of ASM’s control, ASM assumes no liability or obligation in connection with any use of this information. No claim of any kind, whether as to products or information in this publication, and whether or not based on negligence, shall be greater in amount than the purchase price of this product or publication in respect of which damages are claimed. THE REMEDY HEREBY PROVIDED SHALL BE THE EXCLUSIVE AND SOLE REMEDY OF BUYER, AND IN NO EVENT SHALL EITHER PARTY BE LIABLE FOR SPECIAL, INDIRECT OR CONSEQUENTIAL DAMAGES WHETHER OR NOT CAUSED BY OR RESULTING FROM THE NEGLIGENCE OF SUCH PARTY. As with any material, evaluation of the material under end-use conditions prior to specification is essential. Therefore, specific testing under actual conditions is recommended. Nothing contained in this book shall be construed as a grant of any right of manufacture, sale, use, or reproduction, in connection with any method, process, apparatus, product, composition, or system, whether or not covered by letters patent, copyright, or trademark, and nothing contained in this book shall be construed as a defense against any alleged infringement of letters patent, copyright, or trademark, or as a defense against liability for such infringement. Comments, criticisms, and suggestions are invited, and should be forwarded to ASM International. Prepared under the direction of the ASM International Technical Book Committee (2007–2008), Lichun L. Chen, Chair. ASM International staff who worked on this project include Scott Henry, Senior Manager of Product and Service Development; Steven Lampman, Technical Editor; Ann Britton, Editorial Assistant; Bonnie Sanders, Manager of Production; Madrid Tramble, Senior Production Coordinator; Diane Grubbs, Production Coordinator; Patricia Conti, Production Coordinator; and Kathryn Muldoon, Production Assistant Library of Congress Control Number: 2008925435 ISBN-13: 978-0-87170-868-7 ISBN-10: 0-87170-868-X SAN: 204-7586 ASM InternationalÕ Materials Park, OH 44073-0002 www.asminternational.org Printed in the United States of America
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This book is dedicated to our families, without whose continued support the completion of this work would not have been possible: My husband, Antonio Carlos Canale, and my children, Amanda, Sara, and Bruno L.C.F.C. To my lovely wife, Carla Mesquita, and my dear son, Rafael R.A.M. My wife, Alice G.E.T.
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Contents Preface ............................................................................................................................. ix
Component Design ............................................................................................................ 1 Mario Solari, Consultores de Tecnologı´a e Ingenı´era SRL Pablo Bilmes, Universidad Nacional de La Plata Introduction to Heat Treat Processing .................................................................................. 1 Important Design Aspects ..................................................................................................... 2 Techniques for Controlling Distortion ................................................................................ 16 Examples of Failures due to Heat Treatment ...................................................................... 18 Heat Treatment Design ........................................................................................................ 29 Modeling of Heat Treatment ............................................................................................... 31 Failure Aspects of Welded Components ............................................................................. 33 Heat Treatment Procedures Applied to Welded Components ............................................ 36 The Risk-Based Approach and Heat Treatments ................................................................ 40
Overview of the Mechanisms of Failure in Heat Treated Steel Components .................... 43 Scott MacKenzie, Houghton International, Inc. General Sources of Failure .................................................................................................. 43 General Practice Conducting a Failure Analysis ................................................................ 47 Determination of the Fracture Mechanism ......................................................................... 51 Summary .............................................................................................................................. 83
Mechanisms and Causes of Failures in Heat Treated Steel Parts ....................................... 87 Debbie Aliya, Aliya Analytical, Inc. Types of Damage and Failure ............................................................................................. 88 Factors Contributing to Poor Response from Heat Treatment ......................................... 101 Concluding Comments ...................................................................................................... 108
General Aspects of Failure Analysis ............................................................................... 111 Waldek Wladimir Bose-Filho, Universidade de Sa˜o Paulo Jose´ Ricardo Tarpani, Universidade de Sa˜o Paulo Marcelo Tadeu Milan, Instituto de Materiais Tecnolo´gicos do Brasil Ltda. General Guidelines of Failure Analysis ............................................................................ 111 Fracture .............................................................................................................................. 118 Distortion ........................................................................................................................... 127 Wear-Assisted Failure ....................................................................................................... 129 Environmentally Assisted Failure ..................................................................................... 131 v
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Failure in Steel Forging .................................................................................................. 133 Md. Maniruzzaman, Worcester Polytechnic Institute Charlie Gure, Forging Consultant Stephen R. Crosby, The Stanely Works Richard D. Sisson, Jr., Worcester Polytechnic Institute Forging Process Design ..................................................................................................... 134 Case Studies ....................................................................................................................... 138
Failures from the Casting Process .................................................................................. 151 Omar Maluf, Instituto de Materiais Tecnolo´gicos do Brasil Ltda. Luciana Sgarbi Rossino, Instituto de Materiais Tecnolo´gicos do Brasil Ltda. Camilo Bento Carletti, Centro de Caracterizac¸a˜o e Desenvolvimento de Materiais Celso Roberto Ribeiro, Centro de Caracterizac¸a˜o e Desenvolvimento de Materiais Clever Ricardo Chinaglia, Centro de Caracterizac¸a˜o e Desenvolvimento de Materiais Jose´ Eduardo Mya, Centro de Caracterizac¸a˜o e Desenvolvimento de Materiais Failures due to Improper Cast Design ............................................................................... 151 Effects due to Porosity ...................................................................................................... 154 Effects due to Decarburization during Microfusion ......................................................... 162 Effects due to Cold Joints ................................................................................................. 163 Inclusions ........................................................................................................................... 165
Sources of Failures in Carburized and Carbonitrided Components ................................ 177 Malgorzata Przylecka, Poznan´ University of Technology Wojciech Ge˛stwa, Poznan´ University of Technology L.C.F. Canale, Universidade de Sa˜o Paulo Xin Yao, Portland State University G.E. Totten, Associac¸a˜o Instituto Internacional de Cieˆncia and Portland State University Design ................................................................................................................................ 179 Steel Selection and Hardenability ..................................................................................... 181 Residual Stress ................................................................................................................... 196 Dimensional Stability ........................................................................................................ 200 Quenching and Grinding Cracks ....................................................................................... 204 Insufficient Case Hardness and Improper Core Hardness ................................................ 209 Influence of Surface Carbon Content ................................................................................ 211 Influence of Grain Size ...................................................................................................... 217 Internal Oxidation .............................................................................................................. 219 Carbides and Carbide Structure ........................................................................................ 222 Noncarbide Inclusions ....................................................................................................... 228 Micropiting ........................................................................................................................ 230 Contact Fatigue Piting (Macropiting) ............................................................................... 230 Case Crushing .................................................................................................................... 231 Pitting Corrosion ............................................................................................................... 232 Partial Melting ................................................................................................................... 233
Fatigue Fracture of Nitrided Layers ............................................................................... 241 Aleksander Nakonieczny, Institute of Precision Mechanics Fatigue Resistance ............................................................................................................. 241 Fatigue Evaluation of Nitrided Steels ............................................................................... 244 Fatigue Property Characteristics after Carbonitriding ...................................................... 246 Summary ............................................................................................................................ 250 vi
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Steel Heat Treatment Failures due to Quenching ........................................................... 255 L.C.F. Canale, Universidade de Sa˜o Paulo G.E. Totten, Associac¸a˜o Instituto Internacional de Cieˆncia and Portland State University Phase Transformation During Heating and Quenching .................................................... 255 Effect of Materials and Quench Process Design on Distortion ........................................ 263 Stress Raisers and Their Role in Quench Cracking .......................................................... 272 Case Studies in Quench Cracking ..................................................................................... 273 Steel Failures due to Tempering and Isothermal Heat Treatment ................................... 285 Jan Vatavuk, Universidade Mackenzie L.C.F. Canale, Universidade de Sa˜o Paulo Martensite .......................................................................................................................... 285 Tempering .......................................................................................................................... 289 Embrittlement .................................................................................................................... 293 Case Studies ....................................................................................................................... 303 Failure Analysis in Tool Steels ....................................................................................... 311 Rafael A. Mesquita, Villares Metals Celso Antonio Barbosa, Villares Metals Classification of Tool Steels .............................................................................................. 311 Heat Treating Failures of Cold Work Tools ..................................................................... 314 Heat Treating Failures of Hot Work Tools ....................................................................... 330 Conclusion ......................................................................................................................... 349 Case Studies of Steel Component Failures in Aerospace Applications ............................ 351 Scott MacKenzie, Houghton International, Inc. Failure Analysis of a Catapult Holdback Bar ................................................................... 351 Cracking in a Main Landing Gear Attach Pin .................................................................. 354 MLG Linear Actuating Rod and Cylinder ........................................................................ 355 Failure Analysis of AISI 420 Stainless Steel Roll Pin ...................................................... 359 Failure Analysis of a Main Landing Gear Lever .............................................................. 362 Failure Analysis of an Inboard Flap Hinge Bolt ............................................................... 364 Failure Analysis of a Nose Landing Gear Piston Axle ..................................................... 367 Multiple-Leg Aircraft-Handling Sling .............................................................................. 372 Failure Analysis of an Aircraft Hoist Sling during Static Test ......................................... 373 Failure Analysis of an Internal Spur Gear ........................................................................ 375 Main Landing Gear Axle ................................................................................................... 378 Nondestructive Testing and Failure Analysis of Fin Attach Bolts after Full-Scale Fatigue Testing ........................................................ 380 Failure Analysis of Powder Metal Steel Components ..................................................... 395 S. Ashok, Sundram Fasteners Ltd. Sundar Sriram, Sundram Fasteners Ltd. Powder Metallurgy Process ............................................................................................... 395 Case Hardening ................................................................................................................. 397 Failure Analysis Techniques ............................................................................................. 399 Case Studies of PM Steel Failures .................................................................................... 401 Induction Hardening ..................................................................................................... 417 Janez Grum, University of Ljubljana Steels for Surface Hardening ............................................................................................. 419 vii
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Main Features of Induction Heating ................................................................................. Induction Hardening of Machine Parts ............................................................................. Magnetic Flux Concentrators ............................................................................................ Conditions in Induction Heating and Quenching of Machine Parts ................................. Time-Temperature Dependence in Induction Heating ..................................................... Quenching Systems for Induction Hardening ................................................................... Time Variation of Stresses and Residual Stresses ............................................................ Workpiece Distortion in Induction Surface Hardening .................................................... Residual Stresses after Induction Surface Hardening and Finish Grinding ..................... Hardness Profiles in the Induction Surface-Hardened Layer ............................................ Fatigue Strength of Materials ............................................................................................ Stress Profiles in Machine Parts in the Loaded State ........................................................ Input and Output Control of Steel for Induction Surface Hardening of Gears ................
420 422 437 440 444 449 452 466 472 477 481 485 491
Failure Analysis of Steel Welds ...................................................................................... 503 J.H. Devletian, Portland State University D. Van Dyke, MEI-Charlton, Inc. Discontinuities in Steel Welds .......................................................................................... 503 Fatigue of Welded Joints ................................................................................................... 505 Hydrogen-Assisted Cracking Theory ................................................................................ 506 Types of Hydrogen-Assisted Cracking ............................................................................. 509 Stress-Corrosion Cracking of Steel ................................................................................... 513 Solidification Cracking of Steel ........................................................................................ 515 Appendix 1: Metric Conversion Guide .......................................................................... Appendix 2: Temperature Conversion Table .................................................................. Appendix 3: Steel Hardness Conversions ....................................................................... Appendix 4: Austenitizing Temperatures for Steels ........................................................ Appendix 5: Temper Colors for Steels ............................................................................ Appendix 6: Physical Properties of Carbon and Low-Alloy Steels ................................... Appendix 7: AISI to Non-AISI Steel Cross Reference ..................................................... Appendix 8: Non-AISI to AISI Steel Cross Reference ..................................................... Appendix 9: Iron-Carbon Equilibrium Diagram ............................................................. Appendix 10: Isothermal Diagrams of Selected Steels ................................................... Appendix 11: Continuous Cooling Diagrams of Selected Steels .....................................
521 525 529 537 539 541 551 563 585 587 601
Index ............................................................................................................................. 629
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Preface Material failures can lead to many potentially disasterous consequences, including poor product quality, necessary repair or component or equipment replacement, production downtime losses, environmental impact, and even loss of life. Furthermore, failures may arise from not one but various causes, including design, material composition, and, in the case of metals such as steel, improper thermal processing. Therefore, when failures do occur, it is critically necessary to not only identify these failures but also to determine and correct their root cause. This is a primary objective of this work. There are many books, journals, and other references that focus on various aspects of failure analysis. However, there are relatively few that focus on steel failures arising during thermal processing, such as forging, casting, heat treatment, welding, and others. A second objective of this book is to provide a reasonably thorough reference detailing potential failures that may occur during thermal processing and the identification of their root cause, even if it is not specifically the thermal process being considered. An important feature of Failure Analysis of Heat Treated Steel Components is that it not only discusses various causes of a failure and its identification but also integrates this discussion with the metallurgy of the process, thus providing one comprehensive resource. This book was developed as a reference source for use by designers, practicing metallurgists, mechanical and materials engineers, quality-control technicians, and heat treaters. This book also will serve as an important textbook for various advanced undergraduate and graduate courses on either failure analysis or thermal processing of steel. The editors are indebted to the invaluable guidance of many persons in the development and production of this text, including Prof. George Krauss (Colorado School of Mines), George Vander Voort (Buehler Ltd., USA), N. Gopinath and V. Raghunathan (Fluidtherm Technology P. Ltd.), Ross Blackwood (deceased), Larry Jarvis (Tenaxol Inc.), and many others. In addition, the editors are most appreciative of Steve Lampman for his continued patience, guidance, and assistance during the various stages of the preparation of this text. The editors are especially grateful for the support of the chapter authors for the diligence, dedication, and patience involved in their vital contributions to this work. Most of all, the editors are especially appreciative of the support and sacrifices made by their spouses, Antonio Canale, Carla Mesquita, and Alice Totten, without which the preparation of this book would not have been possible. We also express our gratitude to Villares Metals S.A. for their continued and vital assistance and generosity throughout this project. Lauralice C.F. Canale, Ph.D. Sao Carlos, SP, Brazil Rafael Agnelli Mesquita Sumare, SP, Brazil George E. Totten, Ph.D., FASM Seattle, WA, USA
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Component Design Mario Solari, CTI Consultores de Tecnologı´a e Ingenierı´a SRL Pablo Bilmes, Universidad Nacional de La Plata
DESIGN involves different creative aspects: planning, development, procedures, availability, and fitness concerning the materials and processes used to manufacture the component. Design is an iterative process, often based on experience, to provide an assessment of the performance of a component for a certain period of time of expected or intended service life. The design process culminates in a technical specification for the part or system and suitable manufacturing processes. Another obvious aim of design is to prevent failures throughout the component lifetime cycle and avoid situations resulting in severe failure. Heat treating achieves the desired changes in structure and properties, and various types of heat treatments may be employed to meet design requirements for mechanical strength, corrosion, wear, and so on. Heat treatments include stress relieving, austenitizing, normalizing, annealing, quenching, and tempering (Ref 1). Heat treating may also involve chemical or additional physical processes. A systematic procedure for minimizing risks involved in heat treated steel components requires a combination of metallurgical failure analysis and fitness for service with respect to safety and reliability based on risk analysis. The effects of steel heat treatment may include (Ref 1):
Control of microstructure formation Increase of strength, toughness, or perhaps creep resistance Relief of residual stresses and prevention of cracking Control of hardness (and softness) Improvement of machinability Improvement of corrosion resistance or wear resistance
Introduction to Heat Treat Processing Material behavior related to heat treatment can be analyzed by developing models that
involve a complex interrelationship of variables associated with the material, manufacturing processes, and service conditions (Ref 2). The ability of ferrous materials to develop required properties through heat treatment is a broad concept that refers both to the ease with which a material may be heat treated and the resulting in-service fitness of the component. The iron allotropic transformation between more densely packed face-centered cubic iron, nonmagnetic gamma (c) phase designated as austenite, and the less densely packed bodycentered cubic iron, alpha (a) phase designated as ferrite, is the basis for heat treatment of steels. Austenite can dissolve up to approximately 2.0 wt% C and in most steels is not stable at low temperature. On the other hand, the interstitial sites in ferrite are much smaller than in austenite; therefore, ferrite can only dissolve very small concentrations of carbon (0.025 wt% maximum) and is relatively soft and stable at room temperature. The iron-carbon phase diagram shows the compositional limits of the different transformational phases formed by a steel alloy that exist during heating or cooling as a function of temperature. In hypoeutectoid steels (those with 50.80 wt% C), upon cooling two different phases can exist, ferrite and austenite, each containing different amounts of carbon. Upon further cooling, the microstructure of these steels exhibits ferrite grains in a pearlite island. Pearlite is a metastable microstructure formed during austenite decomposition. The pearlite structure is an aggregate consisting of alternating lamellae of ferrite and cementite that is formed on slow cooling during the eutectoid reaction. Cementite is a very hard and brittle compound of iron and carbon (Fe3C). Depending on the thermal history, cementite will appear as lamellae (with ferrite), spheroids, or globules in a ferritic matrix.
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2 / Failure Analysis of Heat Treated Steel Components
Microstructures that are formed upon cooling and the proportions of each are dependent on austenitization time, temperature, cooling history of the particular alloy, and the composition of the alloy. The transformation products from austenite decomposition change from a mixture of ferrite and pearlite to bainite or martensite with increasing cooling rate. Bainite is a two-phase mixture of ferrite and cementite consisting of fine lines of iron carbide in acicular ferrite. Upper bainite has a feathery appearance and forms just below the temperature where fine pearlite is formed. Lower bainite exhibits an acicular microstructure that is formed just above martensite, which is produced at approximately 350 C (660 F). Martensite is a supersaturated solid solution of carbon in alpha iron (ferrite) that is less densely packed than the c body-centered tetragonal lattice and is a magnetic platelike structure formed by a diffusionless shear type of transformation of austenite below the martensite start (Ms) temperature. The amount of transformation depends on the martensitic temperature range (Ms to Mf). (Mf is the martensite finish temperature.) The three forms of martensite are lath, plate, and tempered martensite. Transformation from austenite to martensite results in a volumetric expansion at the Ms temperature. Dimensional changes are possible, depending on the carbon content and microstructural transformation product formed. The volume change (%) is [4.64–0.53 (%C)] for the reaction from austenite to martensite. The two most commonly used transformation diagrams are time-temperature transformation for isothermal transformation, and continuous cooling transformation diagrams. These diagrams can be used to predict steel microstructures and hardness after heat treatment, or they may be used to design a heat treatment process. Heat treating processes include hardening, austenitization, annealing (full annealing, intercritical annealing and subcritical annealing, recrystallization annealing, isothermal annealing, soft annealing, diffusion annealing), normalizing, stress relieving, quenching and tempering, and austempering, and are summarized in Table 1 (Ref 1). Hardening and tempering are common heat treatment processes. If steel is cooled sufficiently fast, without microstructural transformation, thermal stresses can develop. Under
these conditions, the surface of the part is initially cooled much more quickly than the core. Therefore, the specific volume in the core is greater than at the surface, and the reduction in volume at the surface is resisted by the greater volume in the core, resulting in the surface being in tension and the core in compression. After the cooling processes have been completed, the residual-stress distribution between the surface and core is obtained. If the surface stresses exceed the hot yield strength of the material, it plastically deforms, resulting in thermally induced dimensional changes (Ref 3). When steels that undergo transformational changes are quenched, the possibility of the formation of both thermal and transformational stresses must be considered. Steel parts are often tempered by reheating after quench hardening to obtain specific mechanical properties. The tempering process involves heating hardened steel to some temperature below the eutectoid temperature for the purpose of decreasing hardness and increasing ductility and toughness while relieving quench stresses and ensuring dimensional stability. Tempering processes include tempering of martensite, transformation of retained austenite to martensite, tempering of the decomposition products of martensite, and decomposition of retained austenite to martensite. In addition, tempering may also lead to dimensional variation due to relaxation of residual stress and plastic deformation, which is due to the temperature dependence of yield strength. Tempering may lead to an increase in hardness if secondary hardening occurs, which is due to precipitation of a compound or to the formation of martensite or bainite from retained austenite, decomposition during tempering, or destabilization during this process and then transformation during subsequent cooling. Quenchant selection and quenching conditions are critically important parameters in quench system design. For example, the dimensional changes after austenitizing and then quenching in water are greater than quenching in oil (Ref 3).
Important Design Aspects The importance of good design cannot be overemphasized. Poor design can cause or
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Table 1 Heat treating process for carbon steels Process
Characteristics
Austenitization: Complete transformation to austenite by heating the steel above the critical temperature for austenitic formation
The optimal austenitization temperature is 30–50 C (55–90 F) above Ac3 for hypoeutectoid steels and 30–50 C (55–90 F) above Ac1 for hypereutectoid steels. Ac3 is the temperature at which the transformation of ferrite to austenite is completed during heating. Ac1 is the temperature at which austenite begins to form during heating. The heating rate must be limited and uniform to avoid cracking or warpage and to control thermal stresses in the range of 250–600 C (480–1110 F). The carbon equivalent controls the propensity for steel to crack. The holding time is dependent on geometrical factors related to the furnace (emissivities, temperature, and atmosphere composition) and load (type of steel and thermophysical properties).
Annealing: Heat treatment consisting of heating and soaking at suitable temperature followed by cooling under conditions such that, after return to ambient temperature, the metal will be in a structural state closer to that of equilibrium
Full annealing: Heat 30–50 C (55–90 F) above Ac3 for hypoeutectoid steels, then furnace cool through the critical temperature range at a specified cooling rate. The aim is to break the continuous carbide network of high-carbon steels. It improves machinability. Partial (intercritical) annealing: Heating within the critical temperature range (Ac1–Ac3), followed by slow furnace cooling. It improves machinability. Subcritical annealing: Heating 10–20 C (20–35 F) below Ac1 followed by cooling in still air. It can be used to temper bainitic or martensitic structures to produce softened microstructures containing spheroidal carbides in ferrite. Improves the cold working properties of low carbon steels (525% C) or softens high-carbon and alloy steel Recrystallization annealing: Heat the steel for 30 min–1 h at temperature above the recrystallization temperature (TR=0.4 Tm), then the steel is cooled. The treatment temperature depends on prior deformation, grain size, and holding time. The recrystallization process produces strain-free grain nucleation, resulting in a ductile, spheroidized microstructure. Isothermal annealing: Heating the hypoeutectoid steel within the austenitic transformation range above Ac3 for a time sufficient to complete the solution process, yielding a completely austenitic microstructure. At this time, the steel is cooled rapidly at a specific rate within the pearlite transformation range until the complete transformation to ferrite plus pearlite occurs, and then it is cooled rapidly. Spheroidizing (soft annealing): Involves the prolonged heating of steel at a temperature near the lower critical temperature (Ac1), then furnace cooling Diffusion (Homogenizing annealed): Heat the steel rapidly to 1100–1200 C (2010–2190 F) for 8–16 h, furnace cool to 800–850 C (1470–1560 F), and then cool to room temperature in still air. It is performed on steel ingots and castings to minimize chemical segregation.
The primary purpose of annealing is to soften the steel to enhance its workability and machinability. Also, it relieves internal stresses, restores ductility and toughness, refines grains, reduces gaseous content in the steel, and improves homogenization of alloying elements.
Normalizing: The aim is to provide a uniform microstructure of ferrite plus pearlite (small grains and finer lamellae than in annealing).
Heat the steel to 40–50 C (80–90 F) above Ac3 for hypoeutectoid steels and 40–50 C (80–90 F) above Acm for hypereutectoid steels. The holding time depends on the size, and then the steel is cooled in still air. It produces grain refinement and improved homogenization.
Stress relieving: It is typically used to remove residual stresses that have accumulated from prior manufacturing processes. Stress relieving results in a significant reduction of yield strength in addition to reducing the residual stresses to some “safe” value.
Heat to a temperature below Ac1for the required time to achieve the desired reduction in residual stresses, and then the steel is cooled at a rate sufficiently slow to avoid the formation of excessive thermal stresses. Below 300 C (570 F), faster cooling rates can be used. No microstructural changes occur during stress-relief processing. The recommended heating temperature range is 550–700 C (1020–1290 F), depending on the type of steel. These temperatures are above the recrystallization temperature. Little or no stress relief occurs at temperatures 5260 C (500 F), and approximately 90% of the stress is relieved at 540 C (1005 F). The maximum temperature for stress relief is limited to 30 C (55 F) below the tempering temperature used after quenching. The results of the stress-relief process are dependent on the temperature and time.
Hardenability: Ability to develop hardness to a given depth after having been austenitized and quenched
The hardenability depends on the concentration of dissolved carbon in the austenitic phase, alloying elements, austenitizing temperature, austenitic grain size at the moment of quenching, size and shape of the cross section, and quenching conditions.
Quenching: Quench severity is the ability of a quenching medium to extract heat from a hot steel workpiece.
Specific recommendations for quench media selection for use with various steel alloys are provided by standards such as SAE AMS 2759. Quench media include water, brine, aqueous polymer, gas or air quenching, and caustic quenching.
Tempering: Tempering is the thermal treatment of hardened and normalized steels to obtain the desired mechanical properties, which include improved toughness and ductility, lower hardness, and improved dimensional stability.
The tempering process involves heating steel to any temperature below the Ac1 temperature. During tempering, as-quenched martensite is transformed into tempered martensite, which is composed of highly dispersed spheroids of cementite (carbides) dispersed in a soft matrix of ferrite, resulting in reduced hardness and increased toughness. The objective is to allow the hardness to decrease to the desired level and then stop the carbide decomposition by cooling. The extent of the tempering effect is determined by the temperature and time of the process.
Source: Ref 1
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4 / Failure Analysis of Heat Treated Steel Components
promote heat treatment failures before the component is put into service, or it may reduce service life, sometimes dramatically. The objective of proper design for heat treatment is to provide the minimum engineering requirements, the desired material properties at the lowest total cost, and, in particular, to minimize the expense of scrap due to rework of parts that may have undergone excessive distortion or cracked. The Heat-Transfer Theory Applied to Heat Treatments. The laws that govern heat transmission are very important to the engineer in heat treatment design. There are three different types of heat transfer: conduction, convection, and radiation (Ref 4). They have in common that temperature difference (thermal gradient) must exist and that the heat is always transferred in the direction of decreasing temperature. When the temperature profile does not change with time, the fundamental relation for the unidirectional steady flow of heat through a solid by conduction, Fourier’s first law, can be expressed by: Q=7l
¶T ¶x
(Eq 1)
where Q is the quantity of heat flowing through the unit area of a wall per unit time in the direction of the x-axis and is directly proportional to the thermal conductivity, l, and the thermal gradient in x-direction. Thermal conductivity has a nearly linear dependence on temperature. During heat treatments, the temperature varies in time as well as in space; these processes are called unsteady, nonstationary, or transient. As the body heats, the temperature at each point asymptotically approaches the temperature of the medium. The temperature of points near the surface of the body changes most rapidly. The differential equation for one-dimensional transient heat conduction, Fourier’s second law, in the absence of inner heat sources is: ¶T ¶2 T =a 2 ¶t ¶x
(Eq 2)
where T is the temperature, t is the time, and a is the thermal diffusivity of the metal and is: a=
l rCp
(Eq 3)
where l is the thermal conductivity, r is the density, and Cp is the specific heat at constant pressure of the material. The thermophysical
properties l, r, and Cp vary with temperature. The differential equation of heat conduction establishes the relation between the time and space variation of temperature at any point of the body in which conduction takes place. The factor of proportionality thermal diffusivity, a, represents a physical property of the material, is essential for transient processes of heat flow, and defines the rate of change of temperature. If the thermal conductivity, l, is the ability of a solid to conduct heat, thermal diffusivity is the measure of a material thermal inertia. The quantity rCp is the volumetric specific heat; this product is approximately constant for solid metals. So, in the case of austenitic stainless steels, low thermal conductivities correspond to low thermal diffusivities. In other words, equalization of temperature at all points of space will proceed at a lower rate in austenitic stainless steels, with respect to ferritic steel, due to its lower thermal diffusivity, and there are difficulties in homogenizing temperature during heat treatments. Heat transfer by convection occurs between the surface of the body and surrounding fluids; for this type of heat transmission, the following Newton’s equation is in general use: Qh =hDT
(Eq 4)
It simply states that an invariable temperature difference, DT, between a surface and a fluid in contact with it causes a steady heat flow of Qh. The factor of proportionality, h, is called the coefficient of heat transfer. The third type of heat-transfer mode is radiation. The heat flow by radiation is commonly written: Q=es T14 7T04
(Eq 5)
where e is the emissivity (1 for a black body), T is the temperature, and s is the Stefan constant (Ref 5). The three different types of heat transfer, conduction, convection, and radiation, are present during the heat treatment processes. Perhaps the most important physical property of steel to be considered in design is its coefficient of thermal expansion. Most heat treating problems could be solved if this coefficient could be controlled. Because it cannot, it is necessary to learn to design with it (Ref 6). Almost all solids expand on heating. As the temperature is raised, the thermal vibration pushes the atoms apart, increasing their mean
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spacing. The effect is measured by the linear expansion coefficient: a=
1 dl L dT
(Eq 6)
where L is a linear dimension of the body (Ref 5). The relationship between thermal conductivity and thermal expansion is important in designing against thermal distortion. Thermal gradients can cause a change of shape, which is a distortion of the component. The strain is related to temperature by: e=aðT0 7T Þ
(Eq 7)
where T0 is ambient temperature, and a is thermal expansion (Ref 5). The distortion is proportional to the gradient of the strain, so it is proportional to the thermal gradient. By Fourier’s first law, the heat flow is proportional to the thermal gradient through the thermal conductivity, l. For a given geometry and heat flow, the distortion is minimized by selecting materials with large values of l/a (Ref 5). For example, austenitic stainless steels have low thermal conductivity and high thermal expansion, related to ferritic steel, so distortion during welding becomes a problem. Thermal expansion has a strong influence on the development of residual stress. Whenever the thermal expansion or contraction of a body is prevented, thermal stresses appear; if large enough, they cause yielding, fracture, or elastic collapse (buckling). For axial constraint, the stress, Ds, produced by a temperature change of 1 C or the stress per C caused by a sudden change of surface temperature in one that is not constrained is equal to aE, where a is the expansion coefficient, and E is the elastic modulus of the material. For biaxial and triaxial constraint, the stresses shall be multiplied by (1n) and (12n) respectively, where n is Poisson’s ratio. These stresses are large and can cause a material to yield, crack, spall, or buckle (Ref 5). Linear thermal expansion (Table 2) in going from room temperature to 700 C (1300 F) is approximately:
10.3 mm/m (0.124 in./ft) for low-alloy steel 8.5 mm/m (0.102 in./ft) for martensitic stainless steel (type 13Cr) 13.1 mm/m (0.157 in./ft) for austenitic stainless steel (type 18Cr-8Ni).
Thermal expansion of austenite is larger than that of ferrite. These thermal expansion values are representative of solution-annealed material. Subsequent precipitation hardening treatments may affect thermal expansion (Ref 7). In high-temperature components design, incompatibility of thermal expansion becomes a major problem. Choice of materials and designs should take this into account. Bolts used to hold high-temperature casings together must be selected to have sufficient elevated-temperature strength and make a good thermal expansion match with the casing material. When rotor and casing are made of ferritic steel, modified 12% Cr bolts work well up to 565 C (1050 F), but nickel-base superalloys are needed at 595 C (1100 F) or higher. The ability of a material to resist thermal shock, due to a sudden immersion in a cold medium, without cracking depends on its thermal expansion coefficient, a; tensile strength, st, for metals; Young’s modulus, E; thermal conductivity, l; and heat-transfer coefficient, h. A temperature change of DT applied to a constrained body or a sudden change DT of the surface temperature of the unconstrained component induces a stress: s=
EaDT C
(Eq 8)
where C is equal to 1 for axial constraint, (1n) for biaxial constraint, (12n) for triaxial constraint, and n is Poisson’s ratio. If this thermally Table 2 Linear thermal expansion for ferrous materials Linear thermal expansion in temperature range from room temperature to 704 °C (1300 °F) Steel
mm/m
in./ft
Carbon and low-alloy steels: C; C-Mn; C-Si; C-Mn-Si; C-1/2 Mo to 11/4 Cr-1/2 Mo; Mn-1/2 Mo-1/2 Ni 5Cr-1Mo and 29Cr-7Ni-2Mo-N steels 9Cr-1Mo steel 12Cr; 12Cr-1Al; 13Cr; and 13Cr-4Ni steels 15Cr and 17Cr steels 27Cr steels
10.3
0.124
9.5 8.8 8.5
0.114 0.106 0.102
8 7.3
0.096 0.088
13.1
0.157
12.25
0.147
Austenitic stainless steels: 16Cr-12Ni-2Mo; 16Cr-12Ni-2Mo-N; 16Cr-12Ni-2Mo-Ti; 18Cr-8Ni Austenitic stainless steels: 29Ni-20Cr-3Cu-2Mo; 20Cr-18Ni-6Mo; 22Cr-13Ni-5Mn Source: Ref 7
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induced stress exceeds the local tensile strength of the material, yielding (permanent plastic deformation) or cracking results (Ref 5). This plastic flow causes permanent shape change (distortion) and impacts the magnitude and distribution of residual stresses. Water quenching gives a high h, and then the values of DT calculated from the previous equation give an approximate ranking of thermal shock resistance (Ref 5). However, when heat transfer at the surface is poor and the thermal conductivity of the solid is high, the thermal stress is less than that given by the previous equations. A measure of the thermal shock resistance that takes into account the finite rate of heat transfer at the surface, a heat-transfer coefficient that is never infinite, is given by: BDT=
st aE
(Eq 9)
where st, a, and E were defined previously, B=C/A, where C also was defined previously; and A is: A=
sh=l 1+sh=l
(Eq 10)
where s is a typical dimension of the sample in the direction of heat flow, h is the heat-transfer coefficient, and l is the thermal conductivity (Ref 5). The quantity Bi=sh/l is usually called the Biot modulus. If Bi41, heat flow is limited by conduction. For fast water quench of metals, the heat-transfer coefficient, h, is high (h= 104 W/m2K), and the thermal conductivity is also high, so the factor A approaches 1. On the
900
other hand, for fast air flow (h=102 W/m2K), the factor A results are equal to 3 · 10 2 (for section s=10 mm), and the thermal shock resistance DT is larger by the factor 1/A (Ref 5). As an example of the use of the aforementioned equations, Fig. 1 shows schematically the effect of the thermal expansion coefficient (a) and the heat-transfer coefficient (h) in thermal shock resistance (DT) for a hypothetical steel with 800 MPa (~120 ksi) tensile strength, biaxial constraint, and thermophysical properties constant with temperature. Three cases were analyzed: a) In the first case, a ferritic steel with low thermal expansion and a very high-heat transfer coefficient (fast water quench, h=104 W/m2K) was considered. The maximum temperature change (thermal shock resistance) that induces stresses below the tensile strength, avoiding yielding or cracking, is 240 C (470 F). b) The second example similar to case (a) but with high thermal expansion. The thermal shock resistance results in temperatures above 135 C (270 F) increasing the failure risk. c) The third example is similar to case (a) but with a lower heat-transfer coefficient (air flow, h=50 W/m2K). The thermal shock resistance results in 280 C (535 F) decreasing the risk of failure. These factors (residual stresses and dimensional changes) have the greatest influence on
Tensile strength
Stress, MPa (ksi)
800 (120) 700
Induced thermal stress for very high heat-transfer coefficient, h, and low thermal expansion
600 (90) 500
Induced thermal stress for very high heat-transfer coefficient, h, and high thermal expansion
400 (60) 300 200 (30)
Induced stress for low heattransfer coefficient and low thermal expansion
Thermal shock resistance
100 0 0
50 (90)
100 (180)
150 (270)
200 (360)
250 (450)
300 (540)
Temperature change, °C (°F)
Fig. 1
Schematic representation of thermal stresses resulting from a sudden change, DT, of the surface temperature and thermal shock resistance
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the design process of a component. In addition to thermal strains, many materials systems undergo phase transformations as a function of temperature. Often, the new phase(s) that forms has a different volume and different coefficient of expansion as well as different mechanical behavior(s) than the parent phase(s). For example, phase transformation from austenite to martensite results in a volumetric expansion at the martensite start (Ms) temperature. These differences increase the complexity of understanding the effect of thermal gradients on the strains produced and the resulting plastic deformation (Ref 8). Thermal and transformation-induced strains can result in substantial plastic deformation and residual stresses. The total induced strain is the result of the sum of the strain produced by thermal expansion (aE DT=eth) of a piece with initial length (E) and the transformation strain (etr). The total induced strain must be accommodated through either elastic (eel) or plastic (eep) strain, which sums to the total strain et=eth +etr= (eel +eep). In order to determine the accommodation strain values, Young’s modulus (E) and the yield strength (sys) are required as a function of phase and temperature. Most of the plastic deformation occurs during the heat-up and cool-down stages of the process (Ref 8). Primary Stresses, Secondary Stresses, Peak Stresses, and Residual Stresses. Primary stress is a normal or shear stress developed by the imposed loading that is necessary to satisfy the laws of equilibrium of external and internal forces and moments. The basic characteristic of a primary stress is that it is not self-limiting. Primary stress that considerably exceeds the yield strength will result in failure or at least in great distortion. Secondary stress is a normal or shear stress developed by the constraint of adjacent parts or by self-constraint of a structure. The basic characteristic of the secondary stress is that it is self-limiting. An example of secondary stress is a general thermal stress. The elastic stresses calculated previously are nominal values, that do not take into account local discontinuities such as holes, notches, or section changes. Even on a structure where stress intensity has been limited by yield criteria, there may exist highly localized regions where peak stresses are several times higher than yield. Maximum local stresses on a structure can be determined by considering nominal stresses multiplied by a stress-concentration factor and can be estimated through a detailed stress ana-
lysis or by using approximate formulas that account for the most common cases. Design should be verified to confirm whether there are stress-concentration points that may activate failure mechanisms due to brittle fracture, corrosion, or fatigue. Examples of peak stresses are thermal stresses in the austenitic steel cladding of a carbon steel vessel, thermal stresses in the wall of a vessel or pipe caused by rapid change in temperature of the contained fluid, and the stress at a local structural discontinuity. Residual stresses (Ref 9) can be defined as those stresses that remain in a material or body after being manufactured and processed in the absence of external forces or thermal gradients. Residual stresses can be defined as either macro- or microstresses, and both may be present in a component. Macroresidual stresses vary within the body of the component over a much larger range than the grain size. Microresidual stresses, which result from differences within the microstructure of a material, operate at the grain-size level or at the atomic level. Microresidual stresses often result from the presence of different phases or constituents in a material. Residual stresses develop during most manufacturing processes involving material deformation, heat treatment, machining, or processing operations that transform the shape or change the properties of a material. They arise from a number of sources and can be present in the unprocessed raw material, introduced during manufacturing, or can arise from in-service loading. In heat treated parts, residual stresses may be classified as those caused by a thermal gradient alone or a thermal gradient in combination with a microstructural change (phase transformation). When a steel part is quenched from the austenitizing temperature to room temperature, a residual-stress pattern is established due to a combination of a thermal gradient and a local transformation-induced volume expansion. Thermal contraction develops nonuniform thermal (or quenching) stress due to different rates of cooling experienced by the surface and interior of the steel part. Transformational volume expansion induces transformation stress arising from the transformation of austenite into martensite or other transformation product (Ref 10). Residual stresses may be sufficiently large to cause local yielding and plastic deformation,
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both on microscopic and macroscopic levels, and can severely affect component performance. Both the magnitude and distribution of the residual stress can be critical to performance and should be considered in the design of a component. In any free-standing body, stress equilibrium must be maintained, which means that the presence of a tensile residual stress in the component will be balanced by a compressive stress elsewhere in the body. Tensile residual stresses in the surface of a component are generally undesirable, since they can contribute to, and are often the major cause of, fatigue failure, quench cracking, and stress-corrosion cracking. Compressive residual stresses in the surface layers are usually beneficial, since they increase both fatigue strength and resistance to stress-corrosion cracking and increase the bending strength of brittle ceramics and glass. In general, residual stresses are beneficial when they operate in the plane of the applied load and are opposite to it (for example, a compressive residual stress in a component subjected to an applied tensile load). The origins of residual stresses in a component may be classified as mechanical, thermal, and chemical. Mechanically generated residual stresses are often a result of manufacturing processes that produce nonuniform plastic deformation. They may develop naturally during processing or treatment or may be introduced deliberately to develop a particular stress profile in a component. Examples of operations that produce undesirable surface tensile stresses or residual-stress gradients are rod or wire drawing (deep deformation), welding, machining (turning, milling), and grinding (normal or harsh conditions). On a macroscopic level, thermally generated residual stresses are often the consequence of nonuniform heating or cooling operations. These, together with the material constraints in the bulk of a large component, can lead to severe thermal gradients and the development of large internal stresses. An example is the quenching of steel (or aluminum alloys), which leads to surface compressive stresses balanced by tensile stresses in the bulk of the component. Chemically generated stresses can develop due to volume changes associated with chemical reactions, precipitation, or phase transformation. Chemical surface treatments and coatings can lead to the generation of substantial residual-
stress gradients in the surface layers of the component. The criterion applied to avoid plastic deformation states that the calculated stress intensity or effective stress must be lower than the yield and design life creep-rupture stresses of the material. When effective stress is exceeded somewhere within the component, it does not necessarily indicate plastic collapse of the entire structure. Primary stresses may locally exceed yield, within certain limits, provided that there is enough ductility to allow the material to yield without cracking. Plastic collapse occurs when primary stresses are uniform on the entire structure and exceed effective stress. To prevent an incremental collapse or thermal stress ratchet in each loading cycle, the total elastic stressintensity range, considering residual and applied stresses, should be limited to twice the yield stress. Factors Leading to Size and Shape Changes in Heat Treated Components. Within a typical component manufacturing process, there are seven major factors that lead to size and shape changes and the development of residual stresses in heat treated components (Ref 8):
Variation in structure and material composition throughout the component, leading to anisotropy in properties and transformation behavior Movement due to relief of residual stresses from prior machining and forming operations Creep of the part at elevated temperature under its own weight or as a result of fixturing Large differences in section size and asymmetric distribution of material, causing differential heating and cooling during quenching Volume changes caused by phase transformation Nonuniform heat extraction from the part during quenching Thermal expansion
All of these factors, except relief of prior residual stresses (second item) and creep at elevated temperature (third item), can be directly related to thermal and transformationinduced strains in the component (Ref 8). A simple example of how this thermophysical property affects heat treating is given in Fig. 2 (Ref 11). As the shaft is quenched, the corner cools first, and as it shrinks, it mechanically
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upsets the hot steel beneath it. As the quench progresses, the entire shaft cools, but now, because the end is hot upset, the diameter is too small to accommodate the circumference. As a result, the end is (usually) in a high state of residual tensile stress, and if the steel is brittle, quench cracks may develop. The coefficient of expansion is a factor that requires serious design consideration because it affects a part during austenitization. With furnace heating, a part is heated to the austenitization temperature mainly by radiation (80 to 98%) and partly by convection (2 to 20%). By radiantly heating particular portions of a part, thin sections heat fastest, especially those that expose a large surface area, such as a spline or gear. A typical example is the gear on the left in
Fig. 3 (Ref 11). Because its thin sections are relatively rapidly heated, this part could not be made to the required tolerances. The redesign shown on the right was an improvement, but the necessary large access holes were still troublesome. Several other factors at the design stage can contribute to problems traceable to austenitization:
Combinations of components with widely varying (nonuniform) section sizes Designs requiring contact with furnace hearths or placement near walls Designs requiring processing that results in a state of high residual stress before austenitization. Parts that are very thin or long or parts that are large in surface area, which are difficult to heat treat because of distortion during austenitization Designs that are unsuitable for the type of furnace equipment available
In designing a tool or die, various factors must be considered. In practice, it is difficult to separate the design stage from steel grade selection because the two steps are interdependent. The choice of a certain grade of steel, such as one that must be brine or water quenched, will affect all aspects of design and manufacture. In general, any steel grade that requires liquid quenching demands very conservative, careful design. Air-hardening grades tolerate some design and manufacturing aspects that could never be tolerated with a liquid quenching. The design must also be compatible with the equipment available, for example, heat treatment furnaces and surface-finishing
Fig. 2
Effect of coefficient of thermal expansion in heat treating a shaft. Source: Ref 6
Fig. 3
Two gear designs showing the effect of coefficient of thermal expansion. At left is a widely used design, which is very troublesome to heat-treat. A preferred design is shown at right. Source: Ref 11
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devices. Designing tools and dies is more difficult than designing components made from structural steels because of the difficulty in predicting service stresses. Despite advances made in design procedures, much of the design work is still empirically based. Such experience is primarily based on past failures; therefore, it is important that the findings of the failure analyst be incorporated into future work. Despite the shortcomings of the empirical approach, there is a vast body of common-sense engineering
Fig. 4(a, b)
Lathe tool bit of 1.45% C and 1.4% Cr steel with acute-angled and rough-machined crosssectional transition that fractured during hardening. (a) Fracture. Original magnification: 1·. (b) View into angle. 2·. Source: Ref 13
Fig. 5
knowledge available for guidance. Analysis of many tool and die failures shows that two relatively simple design problems cause the most failures. These design shortcomings are the presence of sharp corners and the presence of extreme changes in section mass (Ref 12). A sharp corner concentrates and magnifies applied stresses, stresses that arise in tool and die manufacturing (such as during quenching), or stresses that occur during service. In addition to promoting cracking during liquid quenching, sharp corners promote buildup of residual stresses that may not be fully relieved by tempering and can therefore reduce service life. The largest possible fillet should be used at all sharp corners. Air-quenching grades of steel are more tolerant of sharp corners than liquidquenching grades and are preferred when only minimal fillets can be used. Changes in section size can be the locus of premature failures. Figures 4 to 8 (Ref 13) show failures caused by design errors and selection of unsuitable material. Figure 4(a) shows the fracture of a lathe tool bit made of steel with approximately 1.45% C and 1.4% Cr that was hardened in oil at 870 C (1600 F), which was at least 20 C (35 F) too high. The fracture propagated from a rectangular cross-sectional transition that was not properly filleted and moreover was rough machined, as shown by the grooves in Fig. 4(b) (thereby the notch effect was further aggravated). Many failures in service, especially those caused by shock or cyclic loads, can be caused by such design errors. Figure 5 shows a bolt from a self-service elevator that failed as a result of reverse-bending fatigue. In this case, the fracture also propagated from a sharp-edged cross-sectional transition. To avoid further damage and prevent potential
Bolt of a self-service elevator that failed as a result of reverse-bending fatigue. Source: Ref 13
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Fig. 6
Different parts of rock drills of different durability and made of a steel with 0.95% C, 1.2% Cr, and 0.25% Mo. (a) Broken drills had sharp edges in the hexagonal shaft. Original magnification: 2·. (b) Drills free of defects had well-rounded-off edges. Original magnification: 2·. (c) Fatigue fractures propagated from the sharp edges. Original magnification: 3·. (d,e) Differences are clearly seen in the cross section of the hexagonal shaft. Etching shows that the failed drills also were surface decarburized, which further reduced the fatigue strength. Source: Ref 13
accidents, 24 other bolts that had not yet failed were examined metallographically or in bending tests. Eight of these proved to have incipient fatigue cracks in the cross-sectional transitions. The bolts were partially normalized and partially heat treated. Their strength was determined from Brinell hardness to be between 440 and 700 MPa. Cracks had occurred in the annealed as well as the heat treated bolts, that is, in soft as well as hard bolts. The higher strength of the heat treated bolts was made ineffective by the unfavorable design. Figures 6(a–e) show different parts of rock drills of different durability and made of a steel with composition 0.95% C, 1.2% Cr, and 0.25% Mo. They had failed after a short period of service in the hexagonal shaft, while others had proved free of defects. At first glance, it could be seen that the broken drills had sharp edges (Fig. 6a), while those free of defects were well rounded off (Fig. 6b). Fatigue fractures propagated from the sharp edges. These, in turn, led to catastrophic failures under shock loading (Fig. 6c). The design differences could be clearly seen in the cross section (Fig. 6d,e). Etching showed that the failed drills also were surface decarburized, which further reduced the fatigue strength. Figures 7(a,b) show a compressor transmission shaft with a fracture propagating from an acute-angled keyway, and Fig. 8 shows a drive shaft pinion with fatigue fractures propagating from the acute-angular edge of the helical gear. On the other hand, ignorance, carelessness, and false economies in the selection of materials
cause many errors (Ref 12). Not every steel user is in a position to select the most suitable material for his purpose from the many varieties available. When in doubt, consult the steel manufacturer whose materials specialists possess the necessary knowledge of mechanical and technological properties of the required materials. A close cooperation between the materials specialist of the producer and the designer and plant engineers of the user is the best formula for success. One of the overwhelming causes of steel cracking and unacceptable distortion control is part design (Fig. 9–11). Poor part design promotes distortion, cracking, and nonsymmetrical heat transfer during heating and cooling. Sometimes, designers make designs in which combinations of parts intended to reduce costs can actually increase cost due to problems during austenitization. The classic example is the gear-and-hub combination shown in Fig. 12 (Ref 11). As this part is heated, the thin extremities at the top of the hub heat faster than the sections near the gear. Accordingly, this area has a propensity to increase in size but is restrained by the colder metal nearer the gear; therefore, it upsets itself (yields in compression). Finally, the entire part comes to the prescribed temperature. On cooling, however, and even without quenching, the top end of the hub pinches in because it has upset itself. This upsetting can result in a serious taper condition in the bore. (If the bore is broached before heat treatment, the extremities of the hub will stretch and then close
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in, thus causing additional taper.) The top portion of the hub of the pinion in Fig. 12(a) is used only as a spacer and need not be heat treated. A much shorter hub with a steel tubing spacer (Fig. 12b) would solve the problem in
Fig. 7(a, b)
austenitization, making forging easier and, in most cases, reducing total cost. The bevel pinion shown in Fig. 13 presents a similar problem, although the hub extension is necessary. Here, steel or, preferably, a heat-resistant alloy cap
Compressor transmission shaft with a fracture propagating from the acute-angled keyway. Source: Ref 13
Fig. 10
Design solutions to the distortion problem shown in Fig. 9. Source: Ref 3
Fig. 11
Distortion often encountered when quenching a notch. Source: Ref 3
Fig. 8
Drive shaft pinion with fatigue fractures propagating from the acute-angular edge of the helical gear. Source: Ref 13
Fig. 9
Schematic of a gear that is difficult to harden without the distortion shown. Source: Ref 3
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that will create mass can be put over the thin hub before austenitization to retard the heating rate during carburizing. Because of the extremely rapid heating rate, austenitization with high-frequency electric current can be likened to reverse rapid quenching. Accordingly, design is of utmost importance. As a general rule, steel exposed to the flux of the inductor will heat fastest on corners, around holes (as shown in Fig. 14), and through thin sections. The bottoms of keyways and the roots of gear teeth and splines are austenitized last, often mainly by conduction from adjacent areas.
As in quenching, however, induction tooling can be designed to concentrate flux by using appropriate coil configuration and laminated core material. The different frequencies available provide not only for various depths of heating but also for the sharpness of the heating effect, because the induction-heated layer is often much thinner than the hardened depth of austenitized steel. The extent of conduction is a function of the differential between the surface temperature and that of the core. Thus, preheating, either in the induction coil with a suitable delay or in a furnace, can be employed to reduce heat transfer inward.
Fig. 12
Two designs for gear-and-hub combinations. (a) Difficult to heat treat without excessive taper in the bore. (b) A preferred design. Source: Ref 11
Fig. 13
Redesign of a bevel pinion using electron beam welding that was impossible to heat treat in one piece. Source: Ref 11
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Finally, part design recommendations to avoid distortion and cracking problems (Fig. 15) include:
Parts that exceed the following dimensions often must be straightened or press quenched to maintain dimensional stability: long and thin parts, L=5d for water quenching and L=8d for oil quenching (L is the length, and d is the thickness or diameter); and parts that possess large cross-sectional area (A) and are thin (t), which are defined as A=50t. Balance the areas of mass. Avoid sharp corners and reentrant angles. Avoid sharp corners between heavy and thin sections. Avoid single internal or external keys, keyways, or splines. Provide adequate fillet or radius at the base of gear teeth, splines, and serrations. Do not have holes in direct line with the sharp angles of cutouts. Avoid sharp corners at the bottom of small openings, such as in drawing or piercing dies, because spalling or flaking is likely to result at these points. Keep hubs of gears, cutters, and so on as near the same thickness as possible, because dishing is likely to occur. Order stock large enough to allow for machining to remove decarburized surfaces and surface imperfections, such as laps and seams. Do not drill screw holes closer than 6.35 mm (0.25 in.) from the edges of die blocks or large parts, where possible. Cracking may be avoided by using steel that may be hardened by using lower quench severity, or, if possible, pack the bolt hole to reduce thermal stresses arising due to quenching. Avoid blind holes, if possible. Design all parts with round corners and
Fig. 14
fillets wherever possible. Use air-hardening or high-carbon (oil- and air-hardening) tool steel on unbalanced and intricately shaped dies. Add extra holes, if possible, on heavy, unbalanced sections to allow for faster and more uniform cooling when quenched. Do not machine knife blades to a sharp cutting edge before hardening. Avoid deep scratches and tool marks. The insertion of identification marks on the hardened component is recommended, preferably after hardening, with tools having well-rounded edges and minimum deformation (shallow penetration depth) and at positions far away from the high-stressconcentration zones (reentrant angles, bends, and so on). On long, delicate parallels, shafts, and so on, rough out and have pieces annealed to remove stresses before finish machining. Always use the grade or composition of steel most suitable for the work that the part has to perform. Design symmetry is also an important variable to minimize distortion. A general rule for solving such quench distortion problems is that the “short side is the hot side,” which means that the inside of the bowed metal was quenched more slowly than the opposite side.
Steel Grade and Condition. Although steel cracking is most often due to nonuniform heating and cooling, material problems may be encountered. Some typical material problems include the compositional tolerances: “dirty” steels, those containing greater than 0.05% S (such as SAE 1141 and SAE 1144), are more prone to cracking. It is well known that cracking propensity increases with carbon content. Therefore, the carbon content of the steel is one of the determining factors for
Section through a hole in a part following rapid heating in an induction coil, showing distortion that leads to cracking. Source: Ref 11
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quenchant selection. As a rule of thumb, plain carbon steels with less than 0.35% C rarely crack on hardening, even under severe quenching conditions. The carbon content of steel should never be greater than necessary for the specific application of the part. As a general rule, steels with more than 0.35% C will require oil quenching to avoid cracking. This
Fig. 15
means that higher-cost, higher-alloy steels are required for adequate response to the slower oil quench when carbon content exceeds 0.35%. However, when strength alone (and/or hardness), without the toughness of a quenched and tempered microstructure, fulfills the minimum engineering requirements, the use of
Part design recommendations for minimal internal stresses. Source: Ref 3
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cold-finished bars made with extra-heavy draft or elevated-temperature drawing should receive consideration. In fact, why heat treat when it is unnecessary? Thus, SAE 1050, SAE 1140, or SAE 1144 steels should be particularly attractive to firms with no heat treating facilities or with no commercial heat treater nearby. By adjustment of their composition and the degree to which they are cold (or warm) worked, these steels can be made to have good machining characteristics (Ref 6). On the other hand, it is well known that regions containing high concentrations of coarse carbide microstructure as a result of improper forging may become the initiation point for subsequent quench cracking, particularly with parts of complex shape. It is important to provide a sufficient forging for microstructure to become fine and uniform. Because part manufacture, such as gear production, often requires machining, the condition of the steel that is going to be machined is critically important. Some workers have recommended normalized and subcritical-annealed steels as the ideal condition. The subcritical annealing process reduces the carbon content and alloy carbide content in the austenite, allowing the production of more lath martensite in the microstructure, which provides higher fracture toughness and higher impact toughness. Steel hardenability is determined by its chemistry. The quench conditions required to obtain the desired properties are a function of the hardenability. Therefore, if the steel chemistry is incorrect, the selected quench process conditions may, if too severe, lead to cracking. Unfortunately, this problem is not uncommon.
Techniques for Controlling Distortion In applying one or more of the effective methods of minimizing distortion, cost is usually the major consideration. Therefore, in planning manufacturing operations, it behooves the prudent processor to evaluate the costs of minimizing distortion against the alternatives (Ref 14). In almost any instance, there are at least three alternatives:
Change to another heat treating process Make allowances for stock removal in finishing operations to correct the distortion Incorporate straightening operations as required
In considering the alternatives that relate to minimization of distortion, it is assumed that the grade of material is fixed, and no deviations are allowed in this area. There are often instances, especially for parts of complex design, where a change in steel composition will permit a less drastic quench and thereby reduce distortion. Such changes are usually to steels with higher hardenability. In most instances, however, immediate changes in workpiece composition are not feasible. Pros and cons of the three most likely alternatives (listed previously) are discussed separately in the paragraphs that follow. Consider Change to Another Process. In this area, there are sometimes two or three possibilities, such as changing from a throughhardening steel to a case-hardening type or changing to one that does not require rapid cooling, such as nitriding. One of the most likely changes that is often made in this area is to the use of localized heating, such as induction. For example, a shaftlike member requires hardening only in certain bearing areas. This can be accomplished easily by induction and, in addition to eliminating distortion, is often more economical for other reasons. Parts such as ring gears represent other examples where a change to induction hardening resulted in keeping distortion within acceptable limits. Increase Stock Allowance. In many instances, allowance for stock removal in the finishing operation (usually grinding) is the most economical approach. Under these conditions, some study is usually necessary to determine the magnitude of distortion caused by heat treating and thereby how much stock allowance is required for cleanup. Frequently, it is necessary to take reasonable precautions (perhaps some special procedures) in heat treating and then take further steps by increasing stock removal in finishing. Mechanical straightening, either during processing or applied to heat treated parts, offers a third approach for solving distortion problems. Straightening is sometimes used as the sole technique for correcting distortion, but more often it is used in conjunction with systematic stock removal. Heat Treating Practices for Minimizing Distortion. Positioning in the furnace may have a marked influence on total distortion, especially for parts having a relatively large length-to-cross section ratio. For example, for long, shaftlike parts (solid or tubular), the
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poorest loading technique would be to pile them horizontally and more or less indiscriminately on the furnace hearth. Under these conditions, distortion begins immediately and continues as the parts heat up and lose strength. Parts at the bottom of the pile will naturally distort the most, because they are subject to the greatest stress during heating. In contrast to the poor technique described previously, the best technique for such parts would be to hang them (preferably with spaces between each part) in a vertical furnace for heating. As a rule, some further improvement can be achieved by heating in molten salt as opposed to a gaseous atmosphere. This is due to the fact that some support is supplied by the buoyancy effect of molten salt. One possible disadvantage (relating to distortion) of heating in molten salt is the heating rate. Parts are heated four or five times as fast in a medium of molten salt compared to heating in a gaseous atmosphere. Rapid heating sometimes increases distortion, especially when various section thicknesses are involved. Position during quenching may also have a marked effect on the total amount of distortion. Parts that are hung vertically in the furnace to minimize distortion should likewise be hung vertically in the quenching tank; that is, deep tanks are preferred for this type of work. In most instances, minimum distortion of specific workpieces is achieved when the quenching medium is not agitated. To obtain full hardness, however, agitation usually must be used. When minimum distortion is required, if agitation is used, the quenching medium should be agitated with agitating force at the bottom of the tank. Although this specific system is used for water, the principle applies to any quenching medium. The quenching medium should never be agitated from the side in such a system when minimum distortion is important. Choice of quenching medium can affect distortion. Typical quenching media listed in approximate order of decreasing cooling power, are as follows:
Water Brine solutions (aqueous) Caustic solutions (aqueous) Polymer solutions Oils Molten salts Molten metals Gases, including still or moving
Fog or mists Air
The higher rate of heat extraction (quenching power) is obtained by agitated brine. Minimum distortion would be obtained by vertical heating, then cooling (quenching) by hanging in still air. Of course, this technique would not usually be practical, largely because the parts would have to be made from air-hardening steels. There are several factors that influence choice of quenching medium, but hardenability of the steel is usually the key factor. Cooling rate thus has a marked effect on the amount of distortion. Consequently, the quenching speed should never be faster than is required to attain the required critical cooling rate, when distortion is an important consideration. Special quenching techniques may be needed. Vertical heating and quenching one part at a time could be construed as a special technique. However, commonly recognized special techniques include:
Martempering Press quenching Cold-die quenching
All of these methods can be used effectively to reduce distortion, but as a rule, they are all relatively costly and will greatly increase the total manufacturing cost. Whenever the handling of parts individually is involved, the cost of heat treating increases rapidly. For example, tubes made from 52100 steel, 1.2 m (4 ft) long, 63.5 mm (2.5 in.) in diameter with a 3.175 mm (0.12 in.) wall thickness were required to be hardened to a minimum of 60 HRC with a maximum total indicator reading of 0.75 mm (0.030 in.). No salt bath of sufficient depth was available for quenching; thus, martempering could not be considered. Several different procedures were tried, but the one that was ultimately successful consisted of heating vertically (hanging) and quenching in unagitated oil, one tube at a time. This procedure cost more than six times the estimated cost for bundling the tubes, then mass heating and quenching. Therefore, the cost of such a practice would normally be considered prohibitively expensive. Straightening during processing may have solved the problem, but it also is very expensive. Martempering can, under many conditions, be used to effectively reduce distortion and can still be applied to mass production. This
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depends largely on workpiece shape and size. However, for long, shaftlike or otherwise unwieldy workpieces, martempering requires handling each workpiece individually and is thus a relatively expensive operation, particularly when simultaneous straightening is incorporated. Press quenching is probably the oldest special quenching technique and is still used to a considerable extent. The greatest use of press quenching is for gears that cannot be heat treated with sufficient dimensional accuracy by mass quenching in baskets. Selection of press quenching should be done with the full knowledge that it is very expensive. First, the presses are expensive machine tools. Second, the dies are expensive, and the dies must be tailored to the specific workpiece. Also, press quenching is slow, tedious, and thus expensive. Minimum distortion is achieved by press quenching, but heat treating cost is high. Dry die quenching is another tedious and expensive process and should be considered only for highly specialized applications.
Examples of Failures due to Heat Treatment Different types of errors and failures produced in heat treatment (Ref 15) include:
Heating errors: heating too fast causing stresses in outer zones; heating nonuniformly or locally overheating; heating at too high a temperature or for too long a time (distortion, cracking, residual stresses, decarburization, alloy depletion) Temperature errors: overheating (scaling, burning, internal oxidation, hot shortness, grain coarsening, aging, phases precipitation) Heat treating errors: improper thermal cycle; too high a temperature; too low a temperature; improper heating rate; improper cooling rate; improper soaking from timing errors or nonuniform quenchant; improper atmosphere control, which is critical in carburizing and nitriding; delay between quench and temper; improper aging treatment or postweld heat treatment (improper, unacceptable, or mixed structures and microstructural features; temper embrittlement; sigma-phase embrittlement; 475 C embrittlement; sensitization; carburization;
metal dusting; sulfidation; nitridation; disbonding of chromium-molybdenum steels overlayed by austenitic stainless steels) Influence of design, steel grade, and condition are illustrated in the following examples. Example 1, in Fig. 16 (Ref 16), shows two AISI W1 carbon steel concrete roughers that failed after a few minutes of service. Cracking occurred at the change in section size due to bending stresses. Although the section change has a smooth, filleted surface, it is still a very effective stress concentrator. Subsequent design changes involved a tapered change in section at the cracked location and later at the start of the wrench above the cracked region. Example 2 illustrates that holes placed too close to the edges of components are a common source of failure during heat treatment or in service. Figures 17(a) and (b) (Ref 16) show an AISI O1 tool steel die that cracked during oil quenching. The die face contained numerous fine cracks. The left side of the die broke off during quenching. Figure 17(b) shows both sides of the fracture. Temper color (arrow), typical of the 205 C (400 F) temper used in this case, is apparent. This indicates the depth of the crack produced during quenching that was open during tempering. Coarse machining marks and deep stamp marks were also present. Sharp, unfilleted corners may also promote quench cracking. Example 3, Fig. 18 (Ref 16) shows a 76 by 87 by 64 mm (3 by 37/16 by 2.5 in.) AISI O1 tool steel die that cracked during oil quenching. The cracking pattern (emphasized using magnetic particles) that emanates from the sharp corners is visible. A few cracks are also associated with the holes that are rather close to the edges. Temper color was observed on the crack surfaces, indicating that the cracks were present before tempering. Example 4, in Fig. 19 (Ref 16), shows another example of a quench crack initiated by a sharp corner. This fixture was also made of AISI O1 tool steel that was oil quenched. In this case, the corner was filleted, but there was a nick in the corner where cracking began. The shape of this fixture is also poor for steel that must be oil quenched. The thinner outer regions cool more rapidly, forming martensite first, while the more massive central region cools at a slower rate. An air-hardenable steel would be a better choice for this part.
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Example 5 provides another case of a poor design for liquid quenching, as shown in Fig. 20 (Ref 16). This 76 mm diameter by 76 mm long (3 in. by 3 in.) threaded part made of AISI W2 carbon tool steel cracked in half at an undercut at the base of the threads. Figure 20 shows the two broken halves along with a cold-etched disk taken from the hollow portion of the part. The hardened outer case can be seen in the fracture detail and in the cold-etched disk. Similar parts, without the undercut, were successfully hardened.
Example 6, in Fig. 21 (Ref 16), is a punch made of AISI S7 tool steel that cracked during quenching because of rough machining marks (a common cause of quench cracking). Because of the section size, the punch was oil quenched to 540 C (1005 F), then air cooled. The cracking pattern has been emphasized with magnetic particles. Temper color was observed on the crack walls. Example 7, Fig. 22(a,b) (Ref 16) show a classic example of a failure due to improper
Fig. 17 Fig. 16
AISI W1 (0.85% C) tool steel concrete roughers that failed after short service (2 min for S, 7 min for S11). Failures occur at the change of sections. Source: Ref 16
(a) Front view of an AISI O1 tool steel die that cracked during oil quenching. The die face contains holes that are close to the edge for safe quenching. (b) Side view of broken die halves showing the mating fracture surfaces and temper color (arrow) on the crack surfaces. Source: Ref 16
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Fig. 18
AISI O1 tool steel that cracked during oil quenching. Note the cracks emanating from the sharp corners. The four holes, which are close to the edge, also contribute to cracking. Source: Ref 16
Fig. 19
Fixture made from AISI O1 tool that cracked during oil quenching. The design is poor for liquid quenching. Source: Ref 16
(a)
Fig. 20
electrical discharge machining (EDM) technique. Die cavities are often machined by EDM. The technique has many advantages, but failures have been frequently observed due to failure to remove the as-cast surface region associated with the as-quenched martensitic layer. Cavity surfaces must be stoned or ground, then tempered to prevent such failures. Figure 22(a) shows four 3.2 mm (0.125 in.) diameter EDM holes in an AISI A4 tool steel primer cup plate. The holes were finished by jig-bore grinding, during which spalling was observed at many of the holes (see upper-right hole). The surface was swabbed with 10% aqueous nitric acid to reveal regions affected by EDM. Figure 22(b) shows the microstructure of these regions. An as-cast region was present at the extreme edge (approximately 35.5 HRC). Beneath this layer was a region of as-quenched martensite (approximately 63.5 HRC). Next was a backtempered region (approximately 56 HRC) and then the base-unaffected interior (59 to 61 HRC). The brittle nature of the outer layers and the associated residual-stress pattern caused the spalling. In many EDM-related failures, the as-cast layer is not observed because of the technique used or because of subsequent machining. In these failures, however, an outer layer of brittle as-quenched (white etching) martensite is present. Such a failure is shown in Fig. 23; this failure occurred in a plastic mold die made from AISI S7 tool steel. The crack followed the lower recessed contour of the larger-diameter gear teeth and extended to a depth of approximately 1.6 mm. Etching of the surface revealed an asquenched martensite surface layer (thin, white layer), while the internal structure was grossly
(b) Threaded part made from AISI W2 carbon tool steel that cracked during quenching at an undercut at the base of the threads. Source: Ref 16
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overaustenitized (note the retained austenite, white, and coarse plate martensite). Both factors led to cracking. If the EDM surface layer was not present, poor service life would have resulted anyway due to poor microstructural condition. Example 8, in Fig. 24(a,b) (Ref 16), shows a 41 mm (1.6 in.) square 1.4 kg (3 lb) AISI S5 tool steel sledgehammer head that cracked during quenching. A disk cut from the head was macroetched, revealing a heavily decarburized surface (Fig. 24b). Such a condition promotes quench cracking, particularly in liquid-quenching grades such as S5 (oil quenched), due to differential surface stresses. A deep stamp mark also helped promote cracking. Example 9 presents a situation with stamp marks, such as that shown in Fig. 25 (Ref 16) that commonly promote quench cracks. This was present on an air-quenched die made from AISI S7 tool steel. In this case, the die was not tempered, another prime cause of quench cracking. Example 10, Fig. 26 (Ref 17) shows SAE 4140 grade steel seamless tubing that failed because of quench cracks. During production of hydraulic cylinder housings being fabricated from this steel seamless tubing, magnetic particle inspection indicated the presence of circumferential and longitudinal cracks in a large number of cylinders. Figure 26(a) is a cross section of the tube showing extensive cracking revealed by dye-penetrant inspection. Figure 26(b) is a scanning electron microscope
Fig. 21
Punch made of AISI S7 tool steel that cracked during quenching because of rough machining marks (a common cause of quench cracking). Source: Ref 16
(SEM) micrograph showing intergranular fracture at a crack origin. Figure 26(c) is an SEM micrograph illustrating the brittle mode of failure associated with the fracture. Figure 26(d) is a micrograph showing the typical concentrations of nonmetallic stringers in the tube material, and Fig. 26(e) is a micrograph showing a quench crack with a heavy oxide. Although the steel met the compositional requirements of SAE 4140, the sulfur level was 0.022% and would account for the formation of the sulfide stringers observed. The combination of the clustered, stringer-type inclusions and the quenching conditions was too severe for this component geometry. The result was a high incidence of quench cracks that rendered the parts useless. Example 11 presents the case of six wrist pins designed especially for a high-performance six-cylinder automotive engine (Ref 18) that failed after 4800 km (3000 mi) of normal operation. The wrist pins were made of lowcarbon steel carburized on both the outer and inner diameters. Two failed wrist pins were submitted for examination. Sample 1 had fractured into three pieces (Fig. 27). Sample 2 had not fractured but exhibited circumferential cracks on the surface of the central zone. Some of the cracks had progressed for most of the 360 of the pin. Both samples showed some evidence of scoring on the outer diameter. The fractured faces of sample 1 were battered but showed a fairly smooth annular ring around both the outer and inner diameters, with a ductile and fibrous core. The condition of the fractured faces did not permit the definite establishment of a fatigue failure. Figure 28 shows the dimensions of a pin. The machining on the inside diameter surface (indicated by “B”) was relatively rough. The inner diameter had a raised central section with a small fillet on either side (indicated by “A”). The core (Fig. 29) had a banded microstructure of ferrite and pearlite and contained some MnS inclusions. The case (Fig. 30) showed a tempered martensite matrix with a nearly continuous grain-boundary network of cementite. Other cracks started on the surface caused by the cementite network (Fig. 31). All cracks progressed inward from the carburized surface and circumferentially around the pin. Hardness tests were performed on the cross section of the pin, and the depth of case penetration to a value of 50 HRC was measured. Results indicated a case depth of 0.89 to 1.0 mm (0.035
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×
×
×
×
Fig. 22
×
(a) Surface of an AISI A4 primer cup plate showing spalling at one of the 3.2 mm diameter holes made by electrical discharge machining (EDM) Original magnification: 2.5· . (b) Microstructures associated with the spalled hole in (a) caused by improper EDM technique. Source: Ref 16
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to 0.040 in.). Determination of case depth by visual examination on a microspecimen was 0.89 mm (0.035 in.). The principal causes of failure were inadequate heat treatment of the case and a design that incorporated a raised central section of the inner diameter, which acted as a stress raiser. Rough machining of the inner diameter aggravated the situation. The case with the cementite grainboundary network had not been heated to a high enough temperature or long enough to take the cementite into solution in the austenite. It was suspected that after slow cooling from the carburizing temperature, pins were heated slightly above the Ae1 prior to quenching and then given a low temper. The case was refined, and the core was unrefined. Thus, poor heat treatment, resulting in a brittle grain-boundary network of cementite, and a design that formed locations of stress concentration in the inner diameter were the most probable causes of failure. The pins should be carburized to a double heat treatment to refine both case and core and to eliminate the brittle grain-boundary network of cementite. The pin design should be changed to eliminate the central raised section of the inner diameter to avoid the fillets acting as stress raisers. The machining of the inner diameter should be improved to avoid a rough surface. The depth of the carburized case should be reduced to approximately 0.38 mm (0.015 in.) to increase pin toughness. Example 12 features cracking of an alloy steel bolt. A heat treated, cadmium-plated AISI 8740 steel bolt broke through the head-toshank fillet while being handled during assembly (Ref 19). Dimensions of the alloy steel bolt (MSD 21250-10070) were 15.9 mm (0.625 in.)
(a)
Fig. 23
(b)
diameter, 111 mm (4.375 in.) grip length, and 134.5 mm (5.294 in.) overall length. It was heat treated to a tensile strength of 1240 to 1380 MPa (180 to 200 ksi) and a hardness of 39 to 43 HRC and then cadmium plated per QQ-P-416 type II class 2 (23 h bake). The bolt fractured through the head-to-shank fillet, a type of failure usually traceable to a poorly controlled manufacturing process, such as heat treating (quench cracking) or chemical plating (hydrogen embrittlement). In this instance, delayed cracking caused by hydrogen embrittlement was initially suspected, because the bolt reportedly had passed a magnetic particle inspection. The fracture surface (Fig. 32), has two distinct zones. Zone 1 was covered with a thick layer of baked-on scale. The scale was removed and the area examined using an SEM. The fracture topography shown in Fig. 33 is a combination of tearing and intergranular hairline cracks, features often associated with both hydrogen embrittlement and quench cracking. However, the heavy adherent nature of the scale suggests that it formed on the crack surface at high temperature, that is, during heat treating but before conventional quenching and subsequent cadmium plating. Zone 2 is characterized by equiaxed dimples, a common feature of ductile tension overload (final stage of fracture). The bolt head was cut through its centerline, and the crack cross section was metallographically examined in a zone 1 area to further explore whether failure was caused by hydrogen embrittlement (step cracking, no decarburization) or quench cracking (temper scale). The crack had both heavy decarburization and temper scale (Fig. 34), ruling out hydrogen embrittlement cracking. However, because quench cracking
(c)
Plastic mold die made from AISI S7 tool steel that was found to be cracked before use. A crack followed the lower recessed contour of the large gear teeth and had an average depth of 1.6 mm. Smaller cracks were also observed on the flat surfaces. Source: Ref 16
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occurs at relatively low temperatures, decarburization of the surface could only have occurred if the crack was present prior to heat treating. When asked to clarify the situation, the bolt manufacturer admitted that the part had been “quenched in water from high temperatures to verify dimension integrity and was returned to the production lot, instead of being scrapped.” It was assumed that this uncontrolled quench between hot heading and heat treating caused the bolt to crack. Decarburization and scaling occurred during subsequent heat treating of the cracked part. The crack in the bolt occurred subsequent to the hot heading operation prior to the production run. The bolt was quenched in water, dimensionally inspected, and returned to the production lot instead of being scrapped. The heavy decarburization layer on the crack surface supports this scenario. A schematic of the quench crack formation is shown in Fig. 35. Example 13 involves hydrogen embrittlement failure of several cadmium-plated carbon steel socket head cap screws (Ref 20). The cap screws were part of a slide valve assembly on a regenerator line in a petrochemical plant. The screws were exposed to Gulf Coast atmosphere, with no exposure to a chemical process or significant temperatures. The cap screws failed during initial loading, while maintenance was being performed on the valve. One failed and one unfailed cap screw were sent to a laboratory for analysis. The as-received cap screws (Fig. 36) were visually examined. One of the two screws had fractured at the head-to-shank radius and was missing its head. Both screws had been sectioned in the threaded part of the shank (17th and 19th threads) approximately 25 mm (1.0 in.)
Fig. 24
(a) AISI S5 tool steel hammer head that cracked during heat treatment. The fracture was caused by quench cracking by the decarburized surface (b) and deep stamp mark (arrows). Actual size. Source: Ref 16
Fig. 25
Quench crack promoted by the presence of a deep, sharp stamp mark in a die made of AISI S7 tool steel. Source: Ref 16
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(a)
(b)
(d)
(c)
(e)
Fig. 26
A 4140 grade steel seamless tubing that failed because of quench cracks. (a) Cross section of tube showing extensive cracking revealed by dye-penetrant inspection. (b) SEM micrograph showing intergranular fracture at a crack origin. Original magnification: 90·. (c) SEM micrograph illustrating the brittle mode of failure associated with the fracture. Original magnification: 50·. (d) Micrograph showing the typical concentrations of nonmetallic stringers in the tube material. (e) Micrograph showing a quench crack. Note the intergranular branching and heavy oxide. Original magnification: 400 ·. Source: Ref 17
Fig. 28
Fig. 27
Failed wrist pin (sample 1), showing fractured faces. Source: Ref 18
Schematic of wrist pin. Note stress raisers at “A” and the rough machining on surface “B.” Source: Ref 18
from the head of the screw. A crack was observed in the first thread root below the unthreaded part of the shank in the fractured screw. The screws appeared to have been plated. The crack in the fractured screw was opened to reveal its fracture surfaces. Both the initial and laboratory-opened fracture surfaces of the screw were examined with a stereomicroscope at magnifications of 7 to 45 · and with an SEM at much higher magnifications. The fracture surfaces displayed similar fracture modes. Figure 37 shows the overall fracture surface. Figure 38
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shows a ductile fracture mode, which was observed over the majority of the fracture surfaces. Figure 39 shows an intergranular fracture mode, which was observed around the circumference of the screw, next to the plated surface. The part of the shank containing the fracture surfaces was metallurgically prepared in cross section to look for secondary cracking and pitting. No secondary cracking or pitting was observed.
To determine whether the cap screws were plated, a fracture surface and the outside surface of the unthreaded part of the shank were cleaned of oil and other deposits and analyzed using energy-dispersive x-ray spectroscopy (EDS) in conjunction with SEM examination. The EDS
Fig. 31
Fig. 29
Central longitudinal zone of sample 2, showing banded structure of white ferrite and dark unresolved pearlite with MnS inclusions (light gray). 2% nital etch. Original magnification: 200 ·. Source: Ref 18
Fig. 30
Surface structure along a longitudinal axis of specimen 2. The dark matrix is tempered martensite; the light-colored grain-boundary network is cementite. Nital etch. Original magnification: 200 ·. Source: Ref 18
Macrograph of sample 2 taken along the longitudinal axis, showing cracks emanating from both the inner and outer diameters. Unetched. Original magnification: 15·. Source: Ref 18
Fig. 32 Ref 19
Close-up view of the bolt-shank fracture surface. Note the heavy scale on the zone 1 surface. Source:
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Fig. 33
SEM fractography of a field on the zone 1 surface (see Fig. 32). Note the combination of tearing and intergranular fracture. Source: Ref 19
Fig. 34
Optical micrograph of a portion of the crack along a cross section of the fractured bolt head. Note the decarburization at the surface of the crack. Source: Ref 19
Fig. 35
Schematic of quench crack formation. Source: Ref 19
Fig. 36
As-received socket head cap screws. Arrow indicates a secondary crack in the screw thread root. Source:
Ref 20
Fig. 37
Fracture surface of the crack in the failed screw after the crack was opened in the laboratory. “L” indicates the laboratory-induced overload region. Original magnification: 20·. Source: Ref 20
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Fig. 38
SEM micrograph showing a ductile fracture mode observed over the majority of both fracture surfaces. Original magnification: 1000 ·. Source: Ref 20
Iron
20K
LT= 100 SECS BASE METAL OF SCREW
COUNTS
15K
5000
0
Iron
Manganese
10K
0.000 1.000 2.000 3.000 4.000 5.000 6.000 7.000 8.000 9.000 10.000
(a)
ENERGY
Cadmium
Cadmium Cadmium
5000
Cadmium
COUNTS
10K
Cadmium
15K
keV
LT= 100 SECS PLATING ON THE SCREWS SURFACE
0
0.000 1.000 2.000 3.000 4.000 5.000 6.000 7.000 8.000 9.000 10.000
(b)
ENERGY
keV
Fig. 40
Energy-dispersive x-ray spectroscopy spectra of (a) the base metal of the screw and (b) the plating on the outside surface. Source: Ref 20
results are shown in Fig. 40. It was determined that the screws were plated with cadmium. Superficial Rockwell hardness measurements were taken on the metallographic section. The average hardness for the failed cap screw was
Fig. 39
SEM micrograph showing an intergranular fracture mode, observed around the entire circumference at both fractures in the screw. Structure at top is the base metal; structure at bottom is cadmium plating. Original magnification: 1000 ·. Source: Ref 20
80.5 HRN, which converts to approximately 40 HRC. The presence of a ductile fracture mode at the core and an intergranular fracture mode at the outer surface of a plated bolt is typical of hydrogen embrittlement but could also be stresscorrosion cracking (SCC). However, SCC can be eliminated, because the metallographic results showed no evidence of secondary cracks or other corrosion mechanisms, such as pitting. Many hydrogen embrittlement mechanisms have been proposed, but none is universally accepted. However, the phenomenon of hydrogen embrittlement is widely known. The presence of hydrogen in steel reduces the ductility of the steel and causes premature failure under a static load. The time for failure depends on the stress applied to the component and the amount of hydrogen that has diffused into the steel. A component may fail initially when put under load or may fail several weeks after being loaded. Because of this characteristic, hydrogen embrittlement is sometimes called hydrogeninduced delayed failure. Electroplating is a common cause of hydrogen embrittlement in bolts and screws, because hydrogen is evolved (or liberated) during the process. The screws had been plated, and because no other source of hydrogen was identified, it is likely that the plating process was the source of the hydrogen. After most plating processes, bolts and screws are usually baked for several hours at 190 C (375 F) to diffuse any hydrogen out of the steel. However, when a bolt or screw is cadmium plated, it requires a much longer time (approximately 24 h) for baking, because hydrogen
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diffuses less readily through cadmium than other electrodeposited metals. In the case of the cap screws, the screws either were not baked at all or were not baked for a sufficient period of time and/or at a high enough temperature. The cap screws failed because of hydrogen embrittlement. The most probable root cause was absence of baking or insufficient baking of the cap screws after the cadmium plating process. To eliminate the possibility of future hydrogen embrittlement failures, the screws should be baked at approximately 190 C (375 F) for 24 h.
Heat Treatment Design As was analyzed in previous sections, heat treatments are a series of operations in the course of which a solid ferrous product is totally or partially exposed to thermal cycles to achieve the desired change in structures and properties (Ref 2). The chemical composition of the material may possibly be modified during these operations (thermochemical treatment). In addition to the ability of the heat treatment to achieve desired mechanical properties, heat treatment also produces dimensional changes and residual-stress patterns that, in some cases, can lead to component cracking and distortions. In the following section, a procedure is analyzed to improve the performance of the design process. A typical component manufacturing process includes metalworking, machining, or other forming operations, followed by heat treatment. Different types of heat treatment of steels are usually employed by industry: hardening, austenitization, annealing, normalizing, stress relieving, quenching and tempering, and austempering. Heat treatment processes include component heatup, holding at temperature for through-heat solutionizing, or thermochemical treatments such as carburizing or nitriding, quenching from elevated temperature, postquench tempering, or aging treatment. All the steps can influence dimensional changes, residual-stress patterns, and cracking in heat treated components. The Process of Component Design In order to avoid failures associated with heat treatments, it is necessary to develop a
systematic component design process. The typical phases of component design include planning and requirements definition, concept design, detail design, and test and validation. The component is designed to provide a specific mechanical, thermal, and chemical function throughout its life cycle and is often limited by space, cost, and safety considerations. The selection of materials and manufacturing processes for a cost-effective component design is a complex process and often involves iterative decision making. The iterative nature of design requires a continuous analysis and redesign process. Process design employs stress-analysis tools with stress-concentration factors, design rules based on experimental data, material property databases, and mechanical properties resulting from a broad range of heat treatment processes. Computer modeling is a valuable design tool for heat treated components. A computer process simulation model allows a particular design to be tested under a specific set of process conditions. The computer software can graphically display not only the resulting residual stresses and distortions in the component but also the associated transient evolution of temperature, metallurgical phases, volume changes, and stresses. Regardless of the procedure used for developing the design, at the end of the design process, the required quality of the product should be defined and then described in a technical document, definitive layout, or final design intended to accomplish the product manufacturing. Only the minimum quality needed for the product to perform the function intended should be specified. Overspecifying and including restrictive features in the quality description can cause delays and increase costs to the buyer. The design process requires developing operational definitions. Operational definitions are, for example, a specific test of a piece of material or criterion for judgment. Without operational definitions, a specification is meaningless. A specification for heat treatment containing the clause “Avoid long, thin sections” requires operational meaning of “long, thin section.” This definition depends greatly on the quenching media; however, any section length greater than 15 times the diameter is almost always characterized as such, and the slightest nonuniformity in quench will cause it to distort. As the quenching medium becomes more severe
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(i.e., water, caustic quench), this criterion is reduced to as low as 5 times the diameter. For larger length-to-diameter ratios, consideration should be given to fixture quenching or induction hardening. The component design process can be divided into two phases:
Phase 1, which corresponds to the basic definition of the product (including concept and detail design) Phase 2, which corresponds to the design review aiming to prevent failures and minimize risks
Phase 1 deals with the basic definition of the product. As a first step, the generic type of the material as well as its geometric configuration should be selected. The criterion applied to avoid plastic deformation states that the calculated effective stress must be lower than the yield and design life creep-rupture stresses of the material. If the product to be designed is an element that must withstand not only tension loads but also bending, torsional, and axialcompressive loads, then the combined effect of the applied load type, shape, and size and the material properties should be analyzed. Design for those elements subjected to axial compression should be intended to avoid not only plastic collapse but also elastic instabilities. Elastic instabilities may cause Euler buckling and local buckling. The occurrence of this failure mode depends on the geometry of the elements and on the Young’s modulus of the material. In general, the relationship between the Young’s modulus (E) and the density (r), E/r, should be maximized in order to increase stiffness. By doing so, yield occurs before buckling, whereas by increasing specific strength (sf/r), strength also increases; hence, buckling occurs before yield. During the basic design step, brittle fracture should be avoided. Brittle fracture is associated with very little or no plastic deformation. A material may fail in a catastrophic—brittle—manner under stresses even lower than the allowable design stresses used to avoid ductile failures. The material property that controls brittle fracture strength is toughness. Other factors that have an effect on brittle fracture are material thickness; local stress level, including nominal stresses, residual stresses, and stress-concentration factors; temperature; and loading rate. Carbon and lowalloy steels undergo a transition from ductile failure mode to brittle failure mode at low
temperatures. If the material is heat treated during fabrication, it will have adequate toughness in its final condition. Design must be reviewed in order to minimize the presence of notches and defects that concentrate stresses. When mechanical loads are an alternative, there may be a risk of fatigue failures. In fatigue failures, a crack grows in each loading cycle until the remaining ligament fails due to ductile or brittle fracture. This phenomenon can occur at stress levels lower than the allowable stresses for static loads. It should be emphasized that fatigue failures strongly depend on design and manufacturing quality, which is accomplished by increasing fatigue strength and minimizing stress concentrators. Technical requirements should be complied with at the lowest cost. A phase 1 basic design detailed analysis is beyond the scope of this work. Phase 2 design review has the purpose of assuring that the basic design fulfills the requirements and reviews the design to avoid failures. This step verifies that the basic types of failure modes have been properly controlled by design and determines the types of damages associated with each failure mode in order to implement methods for detection. Material behavior can be analyzed by developing models that relate materials attributes, required functions, and manufacturing processes. Due to the large number of aspects involved, the problem can be simplified by considering blocks of knowledge that correspond to specific mechanisms and functions. Each block of knowledge represents a simplified model that relates some properties to the required functions, through the knowledge provided by materials science and engineering. The use of state-of-the-art criteria, which, in some cases, are based on practical experience, can optimize the accomplishment of the analyzed functions. The results of the analysis are synthesized in the definition of design. Consider two groups of behavior models. The first group relates materials attributes— generally known as properties—that are well defined and individually determined. These properties are component shape and size, modulus of elasticity, Poisson’s ratio, ultimate tensile strength, yield strength in tension, shear strength, compressive strength, ductility, elongation, fracture toughness, hardness, thermal conductivity, thermal diffusivity, thermal expansion, specific heat, density, fracture
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toughness, creep strength, creep-rupture strength, and so on. For these properties, test and criteria have accurate operational definitions. The second group of behavior models includes attributes that involve the complex interrelation of a number of variables associated with materials, manufacturing (heat treatment) processes, and service conditions. For this second group of blocks of knowledge, physical metallurgy is intensely used, together with the laws of mechanics and empiric knowledge.
Modeling of Heat Treatment Modeling of heat treatment processes, like other materials processes such as casting and welding, is quite complex due to the tight coupling of various metallurgical transformations and the associated changes in thermal and mechanical states. A heat-transfer model, coupling with a phase transformation model, a thermomechanical model, and a thermochemical model (Fig. 41), is considered.
Fig. 41
Behavior model to analyze heat treatment design
Heat Transfer Model Heat treatments are a series of operations in the course of which a solid ferrous product is exposed to thermal cycles. There are different types of heat transfer: conduction, convection, and radiation. During heat treatments, the temperature varies in time as well as in space; these processes are called unsteady, nonstationary, or transient. The factor of proportionality thermal diffusivity, a, defines the rate of change of temperature, and the heat-transfer coefficient, h, controls the heat flow through the component surface. The thermophysical properties, l, r, and Cp, vary with the temperature. The geometric design will try to avoid mass asymmetries. Large differences in section size and distribution of material causes differential heating and cooling during heat treating. The thermal cycle induces phase transformation and thermomechanical stresses. Solid-Phase Transformation Model The thermal cycle induces solid-phase transformation. Microstructures that are formed upon cooling and the proportions of each are
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dependent on austenitization time, temperature, cooling history of the particular alloy, and the composition of the alloy. The transformation products from austenite decomposition change from a mixture of ferrite and pearlite to bainite or martensite with increasing cooling rate. Isothermal transformation (time-temperature transformation, or TTT) and continuous cooling transformation (CCT) diagrams can be used to predict steel microstructures and hardness after heat treatment. These diagrams are a set of curves drawn in a semilogarithmic coordinate system with logarithmic time/temperature coordinates that define, in the case of the TTT diagram, for each level of temperature, the beginning and end of the transformation of austenite under isothermal conditions, and, in the case of the CCT diagram, define each variation in temperature as a function of time during cooling, the temperature at which the austenite begins and ends its transformation. Industrial heat treatments consist of heating and soaking at a suitable temperature, followed by cooling at an appropriate rate in order to obtain a structural state closer to that equilibrium (annealing), reduce the internal stresses without substantially modifying the structure (stress relieving), increase hardness by more or less complete transformation of austenite to martensite and possibly bainite (quench harden), or obtain a uniform and fine-grained structure with pearlite (normalizing). Other heat treatments are applied to ferrous products after quench hardening (tempering) or solution treatment (aging) to bring the properties to the required level. The process conditions that shall be taken into account usually include the furnace atmosphere (for example, temperature and carbon potential), heating rates, and quench conditions. Thermomechanical Modeling Primary stress is a normal or shear stress developed by the imposed loading that is necessary to satisfy the laws of equilibrium of external and internal forces and moments. Secondary stress is a normal or shear stress developed by the constraint of adjacent parts or by self-constraint of a structure. The basic characteristic of the secondary stress is that it is selflimiting. An example of secondary stress is a general thermal stress. Thermal stresses are related to temperature by the thermal expansion coefficient. The deformation is proportional to the thermal
gradient through the thermal conductivity, l. Thermal expansion has a strong influence on the development of residual stress; whenever the thermal expansion or contraction of a body is prevented, thermal stresses appear. If large enough, they cause yielding, fracture, or elastic collapse (buckling). The total induced strain is the result of the sum of the strain produced by thermal (aE DT=eth) and transformation (etr) strain due to local transformation-induced volume expansion. The total induced strain must be accommodated through either elastic (eel) or plastic (eep) strain, which sums to the total strain: et =eth +etr = (eel +eep )
(Eq 11)
Thermal Shock. The ability of a material to resist thermal shock, due to a sudden immersion in a cold ambient, without cracking depends on its thermal expansion coefficient, a; the tensile strength, st, for metals; the Young’s modulus, E; the thermal conductivity, l; and the heat-transfer coefficient, h. A temperature change of DT applied to a constrained body or a sudden change, DT, of the surface temperature of the unconstrained component induces a thermal strain. Stress Concentration. The elastic strain and stresses calculated previously are nominal values that do not take into account local discontinuities such as holes, notches, or section changes. Even on a structure where stress intensity has been limited by yield criteria, there may exist highly localized regions where peak stresses are several times higher than yield. Maximum local stresses on a structure can be determined considering nominal stresses multiplied by a stress-concentration factor and can be estimated through a detailed stress analysis or by using approximate formulas that account for the most common cases. Design should be verified to confirm whether there are stress-concentration points that may activate failure mechanisms due to brittle fracture, corrosion, or fatigue. A sharp corner concentrates and magnifies applied stresses, stresses that arise in tool and die manufacturing (such as during quenching), or stresses that occur during service. In addition to promoting cracking during liquid quenching, sharp corners promote buildup of residual stresses that may not be fully relieved by tempering and can therefore reduce service life. The largest possible fillet should be used at all sharp corners. Air-quenching grades of
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steel are more tolerant of sharp corners than liquid-quenching grades. Fatigue fractures propagate from the sharp edges. Residual stresses can be defined as those stresses that remain in a material or body after being manufactured and processed in the absence of external forces or thermal gradients. After the cooling processes have been completed, the residual-stress distribution between the surface and core is obtained. If the surface stresses exceed the hot yield strength of the material, it plastically deforms, resulting in thermally induced dimensional changes and thermal and transformational stresses. In heat treated parts, residual stresses may be classified as those caused by a thermal gradient alone and a thermal gradient in combination with a microstructural change (phase transformation). Distortion is any change in the shape and original dimensions of a ferrous product occurring during heat treatment. For a given geometry and heat flow, the distortion is minimized by selecting materials with large values of l/a. Shape and volume changes during heating and cooling can be attributed to three fundamental causes:
Residual stresses that can cause shape change when they exceed material yield strength Stresses caused by differential expansion due to thermal gradients Volume changes due to transformational phase change
Quench Cracking. If thermally induced stress exceeds the local tensile strength of the material, cracking results. Thermochemical Model Thermochemical treatments may be applied to a ferrous product in the austenite state to obtain a surface enrichment in carbon (carburizing), which is in solid solution in the austenite. Carburizing can be done in gas atmosphere, in solid medium, or in a bath of molten salt. Other thermochemical treatments can be applied to produce surface enrichment in nitrogen (nitriding); in nitrogen and carbon (nitrocarburizing); in sulfur, carbon, and nitrogen (sulfidizing); in silicon (siliconizing); in chromium (chromizing); in boron (boriding); and in aluminum (aluminizing). By error in the heat treatment, the surface can be decarburized. Table 3 shows guidelines to avoid heat treatment failures during the design review (phase 2 design review); the recommendations come from the examples presented in the previous sections.
Failure Aspects of Welded Components Brittle Fracture. Low-temperature/lowtoughness fracture is sudden failure of a structural component that is usually initiated at a crack or defect. This is an unusual occurrence, because design stresses are normally sufficiently low to prevent such an occurrence. However,
Table 3 Phase 2 design review to avoid heat treatment failures Characteristic of the material prior to heat treatment operation
Always use the grade or composition of steel most suitable for the work that the part has to perform. The carbon content of the steel is one of the determining factors for quenchant selection; plain carbon steels with less than 0.35% C rarely crack on hardening. Avoid structural and compositional material heterogeneity. Avoid “dirty” steels (clustered, stringer-type inclusions, high sulfur).
Geometry: shape and dimensions
Avoid asymmetric design. Avoid points of stress concentration (sharp corners, blind holes, reentrant angles, single internal or external keys, deep keyways, splines, holes, grooves, very coarse machining marks, deep scratches, and tool marks). Provide adequate fillet or radius at the base of gear teeth, splines, and serrations. Do not have holes in direct line with the sharp angles of cutouts. Do not drill screw holes closer from edges. Add extra holes, if possible, on heavy, unbalanced sections to allow for faster and more uniform cooling when quenched. Order stock large enough to allow for machining to remove decarburized surfaces and surface imperfections, such as laps and seams, and to correct the distortion.
Heat treatment operation
Verify design suitability for the type of furnace equipment available. Avoid thermal shock. Avoid decarburized surface. Consider contact with furnace hearths or placement near walls. Avoid heat treatment errors. Control positioning in the furnace and during quenching. Incorporate straightening operations as required. Avoid creep of the part at elevated temperature under its own weight or as a result of fixturing.
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some older equipment with thick walls, equipment that may be subjected to low temperature due to an upset, or equipment that may be modified could be susceptible to varying degrees of embrittlement. Brittle fracture is associated with very little or no plastic deformation and with a cleavage fracture or intergranular surface, unlike ductile fracture, which is associated with a fibrous surface. A material may fail in an unstable and catastrophic brittle manner under stresses even lower than the allowable design stresses used to avoid ductile failures. This may occur with a combination of material properties and applied stress levels. The material property that controls brittle fracture strength is toughness. Other factors that have an effect on brittle fracture are material thickness; local stress level, including nominal stresses, residual stresses, and stress-concentration factors; temperature; and loading rate. Carbon and low-alloy steels undergo a transition from ductile failure mode to brittle failure mode at low temperatures. Resistance to crack propagation is measured through fractomechanical tests. Crack propagation will occur when the stress intensity at the crack tip, K, reaches a critical value, Kc (MPa m1/2). Brittle materials are those that remain elastic until breaking (break occurs before yield). According to API 581 (Ref 21) lowtemperature/low-toughness fracture of steel is affected by:
Applied loads: Fracture is less likely at low applied loads. Materials specification: Some materials are manufactured to exhibit good fracture properties or toughness properties. Materials are often qualified for use by performing an impact test that measures the energy needed to break a notched specimen. Fine-grained structures, such as tempered martensite, with low impurity content are associated with a high degree of toughness. Other microstructural elements, such as precipitates, second-phase particles, dislocations, and solutes in a solid solution, contribute to increased yield strength but reduce toughness. Temperature: Many materials (especially ferritic steels) become brittle at a temperature called the transition temperature. Brittle fracture is typically not a concern above 300 C (570 F).
Residual stresses and postweld heat treatment Thickness
Temper embrittlement is one of the main causes of toughness degradation in ferritic steels during high-temperature service. This degradation may lead to component failure during service. The problem arises when some types of steels are exposed to temperatures between 345 and 565 C (650 and 1050 F). Typically, 21/4 Cr-1=2 Mo steels with a bainitic structure are the most susceptible to this phenomenon. Temper embrittlement can also occur in C-1=2 Mo, 1Cr-1=2 Mo, 11=4 Cr-1=2 Mo, 3Cr-1Mo, and 5Cr-1=2 Mo steels. Conversely, 9Cr-1Mo steels are less susceptible. Welded joints (weld metal and heat-affected zone) are the most susceptible zones. In all cases, the solution to the problem lies in alloy purity. Exposures within the critical temperature range may occur during temper or postweld heat treatments or during service, and these conditions should be avoided. However, many components operate within the critical temperature range. The segregation of residual elements antimony, arsenic, phosphorus, and tin toward austenitic grain boundaries is the main cause of temper embrittlement. Also, manganese and silicon play an important role in this segregation, and their content should be limited. Both residual and alloy elements can segregate, but the former can be concentrated up to 300 times their average value in the material. Segregation only occurs in ferrite within a 315 to 540 C (600 to 1005 F) temperature range but never occurs during austenitization. In addition to segregation in grain boundaries, a fine precipitation can occur within the grains, resulting in a strength increase (Mo2C precipitation). The phenomenon associated with changes in grain boundaries causes intergranular brittle fracture. In general, ductility and rupture strength are not affected; nevertheless, both can be reduced under severe conditions. Toughness is affected by up to a 100 C (212 F) shift toward the right of the ductilebrittle transition curve, as evidenced from impact testing. When a material is exposed to a 370 to 565 C (700 to 1050 F) temperature range, property degradation may become irreversible; in this case, temper embrittlement and creep operate
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simultaneously. During equipment operation in a hydrogen service environment, hydrogen may diffuse into the metal. During cooling from the operating temperature (shutdown), the material becomes oversaturated with hydrogen. The combination of thermal stresses and hydrogen oversaturation may lead to hydrogeninduced cracking. If the material toughness has been reduced considerably due to temper embrittlement, the risk of a catastrophic cracking is high. Material degradation due to temper embrittlement may result not only in brittle fractures with catastrophic consequences but also in a reduction of the equipment useful life and in a decrease in the equipment reliability and efficiency, since it may be necessary to operate at lower temperatures to avoid temper embrittlement or to depressurize to avoid stresses when the equipment is cold. Temper embrittlement is reversible. Heat treatment for a short period of time at temperatures above 565 C (1050 F), followed by quick cooling, can restore the initial properties. However, if material thickness is high and cooling rates required by precipitation kinetics are not achieved, embrittlement may reoccur. Moreover, the component may crack during heat treatment due to the effect of thermal stresses. Material thickness over 25 mm (0.98 in.) is more susceptible to brittle fracture due to temper embrittlement. Postweld heat treatments minimize the susceptibility to brittle fracture through this mechanism. Aging Tendency. The tendencies for aging after cold working increase hardness and tensile strength, with a simultaneous reduction in ductility. Therefore, the risk of embrittlement due to welding in cold-worked areas increases. The killed steels show enhanced resistance to aging in normalized conditions. When the degree of deformation is high, the material should be thermally treated (normalized or stress relieved) before the deformed zone is welded. The risk of embrittlement due to welding in cold-worked areas to be welded, without special requirements, is allowable under the following conditions. The relationship between the inner bending radius (r) and the plate thickness (t) shall be: r/ti1.0 and tj4.0 mm (0.16 in.) r/ti1.5 and tj8.0 mm (0.32 in.) r/ti2.0 and tj12.0 mm (0.47 in.)
r/ti3.0 and tj24.0 mm (0.94 in.) r/ti10.0 and all thicknesses
These conditions correspond to an elongation, e=1/(2r/t+1), in the cold-worked areas of 33, 25, 20, 14, and 5%, respectively. Hardening Tendency. The maximum hardness in the heat-affected zone (HAZ) depends on the chemical composition (carbon equivalent) and the cooling rate during welding (t8/5). Carbon most affects hardenability, and its effect and that of other elements have been included in carbon equivalent formulas. The International Institute of Welding carbon equivalent formula, recommended for steel with more than 0.18% C, is: CEIIW =C+
Mn Cr+Mo+V Cu+Ni + + 6 5 15
(in wt%)
(Eq12)
When carbon is 50.18%, it is generally recommended that the percentage of cementite (Pcm) formula be used: Si Mn Cu Ni Cr Mo V + + + + + + 30 20 40 60 20 15 10 +5B (in wt%) (Eq13)
PCM =C+
Hardenability of steel is not necessarily an indicator of HAZ hardness. It is important to control the maximum hardness in the HAZ in order to avoid two main problems: cold cracking (hydrogen-assisted cold cracking) during welding fabrication, in which the limit of maximum hardness is usually 350 HV; and in-service cracking in hydrogen environments, where the maximum hardness shall be less than 220 HV. Maximum HAZ hardness can be accurately calculated as a function of chemical composition and the cooling time from 800 to 500 C (1470 to 930 F) (t8/5) (Ref 22). The HAZ maximum hardness for low-alloy steels can be calculated using the following expressions by Du¨ren (Ref 22). For 100% martensite (alloying elements in wt%): HVM =802 C+305
(Eq 14)
For 100% bainite: HVB = Si Mn Cu Cr Ni Mo V 350 C+ + + + + + + 11 8 9 5 17 6 3 +101
(Eq15)
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For 05%martensite5100: h HVx =2019 C (170:5 log t8=5 ) Si Mn Cu Cr Ni Mo V +0:3 + + + + + + 11 8 9 5 17 6 3 +66 (170:8 log t8=5 )
(Eq16)
where the hardness of martensite is HVM, bainite is HVB, the variable amount of martensite is HVx, and t8/5 is the cooling time between 800 and 500 C (1470 and 930 F). Cold cracking can occur after the welding process if four factors are present: local hydrogen concentration, susceptible weld metal or HAZ, local metal hardness, and a high level of residual stress remaining after welding, which may cause cracking at temperatures less than 100 C (212 F). The causes of cold cracking are related to many factors, including initial weld metal hydrogen content, residual hydrogen content at 100 C, steel carbon equivalent, yield stress of steel or weld metal, heat input, preheat temperature, material thickness, joint restraint intensity, notch concentration factor, welding process thermal efficiency, and others. Cold cracking is a diffusion-controlled phenomenon that requires days or weeks at room temperature to develop cracks. Hydrogen-induced cracking in the HAZ can be parallel to the fusion boundary adjacent to a fillet weld or in the form of toe cracks. Weld metals are by no means immune when the steels possess high yield strength. Thus, cold cracking induces surface-connected cracking or subsurface cracking, which may provide initiation points for further cracking by brittle fracture or fatigue. Segregation Tendency. Element (phosphorus, sulfur, carbon, etc.) segregation impairs weldability. Chemical heterogeneity can contribute to localized increases of hardenability. Thus, a normal chemical composition of the heat of the steel may exhibit hardness higher than the maximum allowable hardness in certain parts of the HAZ, despite a normal chemical composition. From the point of view of segregation behavior, semikilled and killed steels are better than rimmed steels. If segregation zones are involved—as in butt welding—care should be taken to limit penetration and hence minimize weld metal dilution. In addition, suitable filler metal and low-hydrogen basic electrodes should be used. The annealing treatment relieves
internal stresses, restoring ductility and toughness, refining grains, reducing gaseous content in the steel, and improving homogenization of alloying elements. Stress-corrosion cracking (SCC) of austenitic stainless steels can be caused by chlorides and polythionic acids, by hydrosulfuric acid on carbon and low-alloy steels, and by caustic corrosion on carbon steels. Stress-corrosion cracking may arise when a susceptible material is simultaneously combined with certain levels of tensile stresses and a critical environment within a specific temperature range. Tensile residual stresses, resulting from manufacturing processes such as welds, contribute to cause this type of damage. Also, cold plastic deformation causes hazardous residual stresses. Tensile stresses must be reduced by controlling manufacturing processes and design. Stress relieving postweld heat treatments can be used to minimize susceptibility to SCC in austenitic stainless steel, attempting to avoid sensitization by using stabilized or low-carbon grades.
Heat Treatment Procedures Applied to Welded Components Carbon and Low-Alloy Steels Postweld heating (soaking) is an option for carbon-manganese and low-alloy steels. Material of thick-walled low-alloy steel pressure vessels is susceptible to hydrogen-induced cracking, and there are many difficulties in detecting small cracks in the HAZ of heavysection steel by conventional nondestructive examination methods. Therefore, it is very important to reduce to a minimum the risk of hydrogen-induced cold cracking of weldment of thick plate under the heavily restricted condition during the entire welding process. One possible solution is to apply both preheating and postheating (Ref 23). There is a temperature range where weld hydrogen-induced cracking may occur and a temperature above the upper limit where no delay cracking occurs, although the hydrogen content or the restraint intensity is high. Therefore, if postheating is carried out above this critical temperature, the discharge of hydrogen is allowed; then, the material can reach room temperature with a minimum risk of hydrogen-induced cracking. The purpose of soaking is to allow hydrogen diffusion to avoid critical values in the
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weldments. Welding must be performed, maintaining preheating and interpass temperatures that depend on the material, process, thickness, type of joints, and heat input. The postheating must be carried out immediately after welding, not allowing the temperature to be lower than 120 C (250 F). The temperature for postheating generally is 300 C (570 F) for 4 to 6 h, depending on the steel type. Then, the weldment is cooled to room temperature. Generally, the same devices recommended for preheating are used. This treatment does not produce either stress relief or microstructural changes. Postweld heat treatment (stress relieving) may be used for carbon-manganese and lowalloy steels. Postweld heat treatment (PWHT) is a uniform heating of a weldment at a temperature below the critical range to relieve the major part of the residual stresses, followed by uniform cooling in still air. The PWHT is carried out to fabricate vessels to increase fracture toughness and minimize the levels of residual stress, which confers resistance to brittle fractures, tempers material structure, and removes the possibility of SCC. The accrued benefits depend on the material under consideration. It is usually required by codes and customers that manufacturers should heat treat all welds on a thick-wall pressure vessel, including any repair welds. The PWHT is conducted at a temperature and for the period specified in the applicable fabrication codes. There are two possibilities: one is to stress relieve the completed vessel in the furnace as a whole; the other is to stress relieve the subassemblies separately in the furnace and then heat treat the final circumferential seam locally. Techniques for PWHT are somewhat similar to those thermal methods in preheating and soaking. When local treatment is carried out, care must be taken to ensure full through-thickness heating, and similarly, the temperature gradient is such that the length of material on either side of the weld at a temperature exceeding half the treatment temperature is at least 2.5(r t)1/2, where r is the bore radius, and t is the material thickness. Problems can arise where butt welds attaching pipes to nozzles positioned close to nozzle/vessel welds require separate heat treatment. In such cases, it may be necessary to apply individual stress analysis to verify the proposed conditions (Ref 23). For both local and furnace treatment techniques, it is very important to exercise control over
heating rates (particularly at temperatures below 300 C, or 570 F, with complex components), time at soak temperature, soak temperature, and cooling rate to avoid such undesirable events as flame impingement, distortion, overheating, air quenching, and reheat cracking. Very largediameter pressure vessels that have transport difficulties may be erected on site and PWHTed from inside by gas burners, using the vessel as its own furnace (Ref 23). It is well known that the mechanical properties of material are degraded due to stress relieving in some materials (Ref 23). Generally, in low-alloy pressure vessel steels, the yield strength, tensile strength, and toughness diminish, and elongation and reduction of area increase as the temper parameter (TP) is increased: TP=T (log t+20) · 103
(Eq 17)
where T is temperature (K), and t is time (h). The PWHT is not considered to be a significant variable for carbon and carbon-manganese steels up to 50 mm (2 in.) (Ref 23). Postweld Heat Treatment of Stainless Steels Postweld heat treatment is used for stainless steels (Ref 24–26). Welding stainless steels in thick sections, when the thickness exceeds approximately 20 mm (0.8 in.), is a complex operation. In addition to codes or engineering specifications, which may impose definitive procedures, it is important to have sufficient metallurgical background to understand what may happen during welding and the subsequent heat treatment operation. The stresses induced by welding often need to be eliminated if dimensional stability of the construction is to be guaranteed and if resistance to SCC is mandatory. The properties of welded joints of stainless steels in thick sections that may be modified during a PWHT are principally the corrosion resistance and the mechanical properties. The possibility of distortion occurring during heat treatment must also be considered. For the purpose of analyzing PWHT in stainless steels, the materials considered were divided into four groups. Chromium Steels. For both ferritic and martensitic types, some codes recommend preheating in the temperature range from 150 to 400 C (300 to 750 F) to avoid problems with hydrogen-induced cracking in welding these steels.
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Martensitic stainless steels are normally PWHTed between 600 and 800 C (1110 and 1470 F), whereas ferritic steels are sometimes heated between 730 and 800 C (1345 and 1470 F), with rapid cooling in order to avoid embrittlement. Treatments above 900 C (1650 F) in these materials are usually intended as homogenizing treatments in order to achieve better properties after the tempering treatment that follows. The mechanical properties are usually improved after this double heat treatment. Low-temperature stress relieving should not be applied to the straight chromium stainless steels, since it may markedly affect the ductility and toughness. The phenomenon is known as 475 C (885 F) embrittlement, and it is due to the coherent precipitation of chromium-rich ferrite, known as alpha prime, within the miscibility gap of the iron-chromium system. This precipitation leads to a slow increase in hardness accompanied by a corresponding loss of toughness. The alpha-prime phase also decreases the corrosion resistance. This type of structural change can be reversed by an annealing treatment at approximately 600 C (1110 F). In 17% Cr steels containing nickel and molybdenum, the toughness is increased by tempering at 630 to 650 C (1165 to 1200 F), below the temperature where austenite or ferrite is formed. In straight 17% Cr steels, the precipitation of sigma phase can occur between 550 and 800 C (1020 and 1470 F), and it is accompanied by a loss of ductility. The sigma phase is formed only after a very long time and may be eliminated by heat treating above 800 C (1470 F). Soft martensitic stainless steels have resulted in an increasingly worldwide use in petrochemical and chemical plants or industries, gas turbine engines, turbine blades, compressors and discs, and in a variety of aircraft structural and engine applications (Ref 27). They have high proof strength and high toughness even in very low temperatures or thick cross sections (Ref 28, 29). If the 12% Cr stainless steels are used as high-strength structural steels, they must be weldable, formable, and have good impact toughness (Ref 30). Hence, in soft martensitic stainless steels, the carbon content is kept below mass 0.1% to improve weldability by promoting a structure with fewer tendencies for cold cracking, better corrosion resistance, and better toughness. Because of the lower carbon, the addition of 4 to 6% Ni (the most powerful austenite former after carbon and nitrogen) is required to avoid delta ferrite, which is
deleterious to impact toughness. For enhanced corrosion, temper embrittlement, and tempering resistance, 0.5 to 2% Mo is added, depending on the intended use. In order to develop the maximum strength and toughness, the steel must be mostly martensitic after cooling, with limited delta ferrite. The martensite must be tempered to obtain good toughness, ductility, and stresscorrosion resistance. In the as-welded condition, the microstructure consists of low-carbon martensite, some presence of delta ferrite, and retained austenite in agreement with the nickel content of the alloy. Postweld heat treatments are necessary to satisfy the service mechanical property requirements (Ref 31). If high impact values are required, PWHTs such as solution annealing plus tempering or double tempering are necessary (Ref 32, 33). The aim of solution annealing is the homogenization of the microstructure by dissolution of the delta ferrite, which is a nonequilibrium solidification product. The delta ferrite is harmful since it increases the ductile-brittle transition temperature. On the other hand, intercritical tempering at 600 C (1110 F) or double tempering (hypercritical plus intercritical) at 670+600 C (1240+ 1110 F) produces tempered martensite with finely dispersed austenite that is stable and not transformable during cooling (Ref 33). It is known that this austenite, which can be observed only by scanning electron microscopy, increases toughness sharply, although it slightly reduces the strength. It has been argued that when retained austenite is present near a propagating crack, the concentrated strain at the crack tip induces transformation into martensite. This mechanically induced transformation would absorb energy and thus increase the toughness. The associated volumetric expansion of the martensitic transformation would tend to close the crack and relieve stresses at its tip. The latter mechanism absorbs strain energy during fracture and therefore limits crack extension (Ref 34). In many codes for austenitic chromiumnickel stainless steels with low ferrite content and fully austenitic alloys, no PWHT is prescribed. In the case of austenitic chromiumnickel steels, low-temperature treatment (400 to 525 C, or 750 to 975 F) will help to achieve dimensional stability of the construction by reducing peak stresses, although this treatment is not frequently used. Treatment in the temperature range from 550 to 1170 C (1020 to 2140 F) is a true stress-relieving treatment. The
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highest part of the range (960 to 1170 C, or 1760 to 2140 F) involves solution treatment and achieves maximum relief of stress; it may be performed to dissolve most of the carbides and sigma phase or delta ferrite. Treatments in the range from 1040 to 1170 C (1905 to 2140 F), followed by water quenching after annealing, are applied to prevent intercrystalline corrosion and SCC. When PWHT is recommended, it only applies to heavy thickness, and a temperature between 900 and 1000 C (1650 and 1830 F) is chosen, followed by water quenching or air cooling, depending on the thickness of the component. Although it may seem difficult to imagine such treatments being applied to heavy components, there are some industrial examples. The stress relieving of austenitic steels is not usually applied except for very thick sections. An exception is in the case of cladding, where the stress-relieving temperature is chosen with respect to the base material, and it is frequently in the range of 540 to 700 C (1005 to 1290 F). Austenitic stainless steels have a higher coefficient of thermal expansion and lower thermal conductivity than ordinary ferritic steels, so greater distortion of welded components must be expected. Unstabilized Austenitic Stainless Steels (UNS S30400, S31600, S30403, and S31603). These grades normally possess excellent weldability, provided they are welded with filler metals that yield an austenitic-ferritic weld metal to avoid hot cracking during welding (5 to 15 ferrite number). It is necessary to follow certain procedures in order to achieve sufficient corrosion resistance, cracking resistance, and toughness. It is well known that austenitic stainless steels are always subjected by the steel manufacturer to a solution-annealing treatment, normally in the range of 1050 to 1100 C (1920 to 2010 F). In the course of this heat treatment, carbide M23C6, sigma phase, and delta ferrite are completely dissolved, and the annealing process produces a homogeneous, fully austenitic structure. With a subsequent quenching treatment, this state is maintained up to room temperature. If possible, the PWHT of welded components should be avoided, with the exception of a solution-annealing treatment. However, if heat treatment cannot be avoided, special attention must be paid to the influence of carbide and phase precipitations on the corrosion-resistance and toughness properties of the weld. The precipitation of sigma phase is the most important of all precipitation phenomena,
apart from M23C6 precipitation, particularly regarding mechanical properties. Due to the highest chromium content of delta ferrite, the weld metal containing delta ferrite is often more precipitation prone than the base metal of similar composition (because the solution-annealing treatment normally applied to the base metal dissolves any delta ferrite). Stabilized Austenitic Stainless Steels. This group of steels contains grades that are alloyed with titanium or niobium in order to improve their intergranular corrosion resistance. Titanium- and niobium-stabilized steels can be welded using niobium-stabilized filler metals with delta ferrite contents in the range of 7 to 15 ferrite number. Whenever possible, heat treatment after welding, for example, stress-relieving treatment, should be avoided. If it cannot be avoided, it is important to use special filler metals with less delta ferrite content to avoid sigma precipitation and detrimental effects on intergranular corrosion resistance. Stabilized steels are somewhat more susceptible to sigmaphase precipitation or to knife-line attack. Fully Austenitic Stainless Steels. This group of stainless steels has a stable austenitic structure that must normally be welded with fully austenitic filler metals that do not produce any ferrite in the weld deposit. In the event of a PWHT, the intergranular corrosion attack range is relatively strongly influenced by the chromium and nitrogen content. The presence of a small amount of sigma-phase or chi-phase precipitation is sufficient to give a marked drop in the pitting resistance. A 475 C (890 F) embrittlement does not occur in fully austenitic weld metal, because this only occurs in the presence of a ferritic structure. Duplex austenitic-ferritic chromiumnickel stainless steels having delta content in the range of 30 to 60% are considered. An alloy containing 25% Cr and 5% Ni, which, in metallurgical terms, is very close to the widely used duplex stainless steel 22Cr-3Mo-5Ni, solidifies completely to delta ferrite from melting. During further cooling, d-c transformation starts at approximately 1200 C (2190 F) with the precipitation of predominantly nodular austenite at the ferrite grain boundary. During further cooling to room temperature, there is only a partial transformation to austenite. The structure now contains some 60% primary delta ferrite and approximately 40% secondary precipitated austenite. If such an alloy is now subjected to a solution-annealing treatment, this ratio can be
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shifted to slightly higher austenite contents. Preheating of duplex steels is not normally required. With thick material, a preheating between 100 and 150 C (212 and 300 F) may be advantageous. In order to obtain high ductility in the welded joint, a solution-annealing treatment, followed by water cooling of the completed welded component, is not normally necessary. If it is required, however, the temperature must be set according to manufacturer’s specifications. With a solution-annealing treatment in the range of 1020 to 1100 C (1870 to 2010 F), very close to a metallurgical equilibrium, it is possible to reverse any harmful structures in the HAZ that may have occurred during welding. It should be noted that this may lead to severe distortion. Heat Treatment of Austenitic-Ferritic Dissimilar Joints. The PWHT of austeniticferritic dissimilar joints or weldments should be avoided whenever possible. However, heat treatments such as annealing or stress relieving may at times be unavoidable or even mandatory. They should always be adapted to suit the requirements of the low-alloy steel section of the joint. Often, the steel in question is a lowalloy creep-resistant steel type used in boiler and pressure vessels that demand a PWHT to suit each particular steel grade. In austeniticferritic dissimilar joints, such heat treatments may lead to the occurrence of the following phenomena:
Carbon enrichment in the weld metal due to the diffusion of carbon from the low-alloy steel into the austenitic weld metal As a consequence thereof, carbon depletion in the HAZ of the low-alloy steel Coarse grain formation in the HAZ of the low-alloy steel due to recrystallization processes Embrittlement of the austenitic weld metal due to precipitation of brittle phases, for example, sigma phase
All these processes are time and temperature dependent. Considering that the normal annealing time is in the range of 2 to 10 h, eventual damage will normally occur, depending on the material combination, at temperatures above 600 C (1110 F). Nickel-base weld metals, due to their high nickel content outside the range of sigma-phase precipitation, show no signs of embrittlement during heat treatments. The use of nickel-base filler metals is recommended if austenitic-ferritic dissimilar joints are subject
to PWHT at temperatures above 600 C (1110 F). Nickel alloy filler metals, abundant and versatile, are frequently used for piping and pressure vessel applications in refineries, chemical plants, and power plants (Ref 35).
The Risk-Based Approach and Heat Treatments Risk analysis is a powerful tool to rationalize the decision-making process and is applicable to heat treated steel components. Increasing incidents led to regulatory action. Some regulations require that recognized and generally accepted good engineering practice (RAGAGEP) must be followed. This also created the desirable objective for industry to document what RAGAGEP was. The concepts of risk-based design, inspection, and maintenance have been developed and are being implemented (Fig. 42). Risk is the combination of the probability (or frequency of occurrence) and consequence (or severity) of a hazard. Its scope is limited to a specific environment during a certain period of time. The intent of risk-based initiatives is to use finite resources and allocate these resources in a manner that achieves the greatest overall reduction in risk. Flaws that exceed the limits permitted by codes may be found. After a defect that is not acceptable is found via a flaw evaluation/fitness-for-service analysis, a repair or replacement is required for continued operation. Various types of damage may occur throughout the different heat treatment stages. These types of damage may lead to failures in the product during its useful life. The term failure refers to the inability of part of the component or the entire component to perform the functions it was designed for, leading to an unprofitable and technically useless product. Failure modes analysis is a procedure in which each potential failure mode is analyzed to determine its effects and the criticality of these effects on the system and to rank each potential failure according to its severity. The criticality analysis involves the use of risk analysis techniques based on the assessment of the likelihood of failure and its potential consequences. Among the most widely used tools in the analysis of design is failure mode and effects analysis (FMEA), which is accepted by the Occupational Safety and Health Administration
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Component Design / 41
Risk Based Initiatives Economic Factor
Risk-Based Design
Risk-Based Inspection Risk Assessment, HAZOP, FMEA, WI, FTA, RCA, RCM, RBI, Fitness-For-Service, Life Extension Material Science FEA, Welding, etc.
FitnessFor-Service
Failure Analysis
Human Factor
Mechanical Integrity Initiatives
Fig. 42
Strategies to minimize risk throughout the cycle life based on risk and mechanical integrity initiatives
(OSHA). These methodologies are based on reliability and safety. By using the FMEA methodology in the analysis of a new design, it is possible to identify single-point failures (a single-point failure refers to an individual failure that may cause the entire system to collapse) and redesign the product to avoid them, thus eliminating them completely or achieving a more robust redesign that is less sensitive to failures. The FMEA has evolved from an ad hoc technique, dependent on a designer’s experience, to a formal and accepted analysis technique. Failure modes can be eliminated by removing their causes or at least having their probabilities of failure reduced to acceptable levels. There are several methodologies to assist in logical thinking to resolve undesirable events. These tools include, among others, reliabilitycentered maintenance, risk-based inspection, FMEA, modified FMEA, and root-cause analysis. Working into an integrated system with riskbased efforts, fractomechanical-based structural integrity approaches, and failure analysis will
allow the development of new tools to assist the designers and manufacturers in minimizing failures related to heat treating operations.
REFERENCES
1. L. Campos Franeschini Canale, G. Totten, and D. Pye, Heat-Treating Process Design, Handbook of Metallurgical Process Design, G. Totten, K. Funatani, and L. Xie, Ed., Marcel Dekker Inc., New York, 2004 2. M. Solari, Risk Based Design, Chapter 2, Handbook of Mechanical Alloy Design, G. Totten, K. Funatani, and L. Xie, Ed., Marcel Dekker Inc., New York, 2003 3. G. Totten, M. Narazaki, R.R. Blackwood, and L.M. Jarvis, Failures Related to Heat Treating Operations, Failure Analysis and Prevention, Vol 11, ASM Handbook, ASM International, 2002, p 192–223 4. M. Jacob and G.A. Hawkins, Elements of Heat Transfer, John Wiley & Sons, Inc., 1957
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5. M.F. Ashby, Material Selection in Mechanical Design, Butterworth Heinemann, 1999 6. R.F. Kern and M.E. Suess, Steel Selection, John Wiley & Sons, 1979, p 35 7. “ASME Boiler and Pressure Vessel Code, Section 2 D,” ASME International, 1999 8. W.E. Dowling, Jr. and N. Palle, Design for Heat Treatment, Materials Selection and Design, Vol 20, ASM Handbook, ASM International, 1997 9. F.A. Kandil, J.D. Lord, A.T. Fry, and P.V. Grant, A Review of Residual Stress Measurement Methods—A Guide to Technique Selection, Report MATC(A)04 Project CPM4.5, Measurement of Residual Stress in Components, Materials Centre, Middlesex, U.K., Feb 2001 10. R.W.K. Honeycombe and H.K.D.H. Bhadeshia, Steels: Microstructure and Properties, 2nd ed., Arnold, 1995 11. R.F. Kern, Selecting Steels and Designing Parts for Heat Treatment, American Society for Metals, 1969 12. K.E. Thelning, Steel and Heat Treatment, Butterworth, London, The Institute of Material, 1993, p 637 13. F.R. Hutchings and P.M. Unterweiser, Failure Analysis, American Society for Metals, 1981, p 35–42 14. H.E. Boyer, Quenching and Control of Distortion, ASM International, 1988, p 245 15. Failure Analysis and Prevention, Vol 11, ASM Handbook, ASM International, 2002 16. G. Vander Voort, Failure Analysis and Prevention, Vol 11, ASM Handbook, ASM International, 1995, p 565 17. Failure Analysis and Prevention, Vol 11, Metals Handbook, 9th ed., American Society for Metals, 1986, p 335 18. W.B.F. Mackay, Failure of Wrist Pins in an Automotive Engine, Handbook of Case Histories in Failure Analysis, Vol 1, ASM International, 1992 19. E. Levy, Cracking of an Alloy Steel Bolt, Handbook of Case Histories in Failure Analysis, ASM International, 1992 20. G.M. Tanner, Hydrogen Embrittlement Failure of Socket Head Cap Screws, Handbook of Case Histories in Failure Analysis, ASM International, 1992, p 332
21. Risk-Based Inspection, 1st ed., API Publication 581, American Petroleum Institute, Washington, D.C., 2000 22. C. Du¨ren, “Equations for the Prediction of Cold Cracking in Field-Welding Large Diameter Pipes,” IIW Document IX-135685, The International Institute of Welding, Cambridge, England, 1985 23. R.W. Nichols, Ed., Developments in Pressure Vessel Technology 3, Applied Science Publishers LTD, London, 1980 24. E. Folkhard, Welding Metallurgy of Stainless Steels, Springer-Verlag, New York, 1994 25. A.W. Marshall, Document IIW-IX-H-42298, The International Institute of Welding, Cambridge, England, 1998 26. J.C.M. Farrar, Document IIW-IX-H-42398, The International Institute of Welding, Cambridge, England, 1998 27. P. Bilmes, C. Llorente, and M. Solari, Role of the Retained Austenite on the Mechanical Properties of 13Cr-4NiMo Weld Metals, Proc. the 20th ASM Heat Treating Society International Conference and Exposition, Oct 2000 (St. Louis, MO), ASM International 28. P. Brezina, Escher Wyss News, Vol 1–2, 1980, p 218 29. H. Niederau, Stahl Eisen, Vol 98 (No. 8), 1978, p 385 30. F.B. Pickering, Physical Metallurgy and the Design of Steels, Applied Science Publishers, London, 1978, p 165 31. T.G. Gooch, Weld. J., July 1995, p 213s 32. R.D. Kane, Corrosion, Vol 33 (No. 7), 1977, p 231 33. P. Bilmes, C. Llorente, and M. Solari, Effect of Post Weld Heat Treatments on the Microstructure and Mechanical Behaviour of 13Cr-4NiMoL and 13Cr-6NiMoL Weld Metals, Proc. the 18th ASM Heat Treating Society Conference and Exposition, Oct 1998 (Chicago, IL), ASM International 34. P. Bilmes, C. Llorente, and M. Solari, Characteristics and Effects of Retained Austenite Resulting from Tempering of 13Cr-NiMo Weld Metals, Mater. Charact., Vol 46 (No. 4), April 2001, p 285–296 35. H. Ebert, Nickel Alloy Filler Metal Review, Weld. J., July 2004, p 60–62
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Failure Analysis of Heat Treated Steel Components L.C.F. Canale, R.A. Mesquita, and G.E. Totten, editors, p 43-86 DOI: 10.1361/faht2008p043
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Overview of the Mechanisms of Failure in Heat Treated Steel Components Scott MacKenzie, Houghton International, Inc. “Primum non nocere” — “First do no harm,” attributed to the ancient Roman physician Galen. “Declare the past, diagnose the present, foretell the future; practice these acts . . . make a habit of two things — to help, or at least to do no harm” (Ref 1).
FAILURES IN STEEL components, like any other material, may have various consequences, such as:
Making the device or component completely inoperable Preventing an operable device from functioning satisfactorily Making the device or component unsafe or unreliable, with immediate removal from service required
Many aspects may also be involved in tracing back to the possible sources of failure of a component. Some of these sources include:
Design Material issues, such as improper materials selection or material imperfections (laps, seams, inclusions, porosity, etc.) Fabrication and processing Rework Assembly Inspection Storage and shipment Service conditions Maintenance Unanticipated service conditions
Many times, more than one factor contributes to a part failure. Rarely is it only one factor.
General Sources of Failure Design deficiencies are a common source of component failure. Examples include the presence of a sharp notch in regions of high stress or a fillet radii that is too sharp. Using a component design for a new application can also lead to
unanticipated failures. Higher stresses or unanticipated service conditions can cause unforeseen failure because of complex or increased stress fields. Stress concentrations may become more critical because of the increase in loading for the new application. Insufficient design criteria can also be the cause of unforeseen failures. Inadequate knowledge of the stress state in the component or inadequate stress calculations can contribute to failures. Much higher stress states than initially assumed or improper stress assumptions can result in premature service failures. Lack of consideration of severe environmental, fatigue, or impact conditions may contribute to failure. Material issues can usually be attributed to either selection of material or material imperfections rendering it unsuitable for service. Inadequate material data can also result in conditions that may contribute to failure. For example, adequate fatigue data, elevated-temperature tensile data, or creep or corrosion data may not be available, and the designer may have to extrapolate or estimate the effects or these properties. Other sources of failure can be attributed to material imperfections. For wrought products, this could be related to segregation, inclusions, porosity, laps, and seams. For castings, these imperfections could be cold shuts, inclusions, shrinkage, voids, and porosity. Forgings can have laps, seams, segregation, and anisotropy in properties from forging flow lines. In one example (Fig. 1), a large roll was heat treated, and several large cracks were observed after inspection. This was originally attributed to quench cracking. On further examination, it was determined that a lap was present in the forging,
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indicated by the presence of high-temperature oxides in the crack along the crack faces. Manufacture and Processing. Processing can have a large influence on properties and the resulting residual stresses. Typically, this is related to wrong procedures or improperly specified procedures. Ambiguous processes or specifications can also contribute to failures due to interpretation or application. Simple things like improper selection of processing sequences or procedures or specifications that were not followed can also contribute to failure. Cold forming, such as stretching or deep drawing, can develop highly localized residual stresses. Local changes in microstructure can occur. Because of the changes in reduction, a large anisotropy in material properties also results. Due to the drawing operation, cracks or microcracking can occur. This could be due to improper lubrication or improper die design. The localized changes in ductility can also contribute to failure. Machining and grinding can create high residual stresses from either machining practice (feeds and speeds) or improper cutting tool selection, material, or geometry. Grinding, if
Fig. 1
abusive, can cause large temperature gradients and localized overheating. This overheating can cause changes in microstructure—either localized softening of the material or localized transformation to martensite and other transformation products—resulting in hard spots. In Fig. 2, a large gear was ground after heat treatment. Because of abusive grinding, local temperatures exceeded the austenitization temperature, and transformation to martensite occurred upon cooling. This transformation and the resulting residual stresses caused cracking of the gear. Temper etch examination of the gear using dilute nitric acid in water in the regions of cracking showed evidence of localized abusive grinding. Identification of parts can also cause failure to initiate. This is from localized impact or electroetching. Localized mechanical stress concentrations or changes in microstructure can occur. This creates either a mechanical or microstructural notch or stress concentration. Heat treatment can cause a variety of different root causes for failures. Overheating, decarburization, quenching, tempering, annealing, and other heat treatments can cause failure to occur.
A large roll was found to have cracks on the outer and inner surfaces of the forging. These cracks were found during final inspection. During examination of metallographic sections taken from the roll, high-temperature oxides were found on the crack faces, strongly suggesting forging laps.
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This could also include improper austenitization temperatures and times. Decarburization is the result of a low-carbon surface from improper atmosphere control. Typically, there is a depleted carbon layer at the surface that, when quenched, is softer than the core material. This soft layer can be completely devoid of carbon (complete decarburization) or only partially depleted in carbon (partial decarburization). This decarburized layer can contribute to premature fatigue failures, because the surface material is different than the designer expected, or failure can result from high residual stresses created at the surface from the quenching operation. The low-carbon surface area can also result in distortion—again, high residual tensile stresses at the surface with low surface hardness. Carburization is similar to the effects of decarburization. In this case, there is a higher surface carbon than expected. High residual tensile stresses can result as well as increased distortion. Quenching can also contribute to high residual stresses or the formation of cracks or microcracking. Transformation stresses from quenching cause the high residual stresses. These high residual tensile stresses can drastically reduce the fatigue strength or have other ramifications in service. Overheating can cause excessive grain growth, with resulting increases in hardenability and increased embrittlement. Underheating can cause poor mechanical properties, because there was an incomplete transformation to austenite and therefore an incomplete transformation to martensite. Poor mechanical properties, such as low tensile and yield stress, and poor impact properties may occur.
Fig. 2
There are also several embrittlement mechanisms caused by the use of improper tempering temperatures. Temper embrittlement and blue brittleness are just two of the common mechanisms that can occur from improper heat treatment and tempering operations. Cleaning, pickling, and electroplating operations can also cause potential failures or contribute to them. Hydrogen charging of highstrength steels from the dissociation of hydrogen on the surface of high-strength steel can occur from cleaning operations in acids. Charging of hydrogen from high current densities in electroplating can cause hydrogen embrittlement unless proper baking procedures are used to allow the hydrogen to diffuse out. Electroplating can also cause high residual tensile stresses, which can contribute to crack initiation. Welding can cause many different problems. These problems can be cracks that are initiated from improper welding procedures, high residual stresses, porosity from inadequately dried weld rods, or dirty workpieces. Microstructural notches or stress concentrations from the heataffected zone and the transition to the base material can be the result of improper preheat and postheat. Improper weld penetration, weld geometry, and excessive weld current (undercutting) can also cause mechanical stress concentrations (Fig. 3). The mast arm failure shown in Fig. 3 (Ref 2) was the result of weld bead undercutting and poor weldment design. Fatigue cracking initiated at the site of the weld toe undercut. This location was a highly stressed area and the location of a large mechanical stress-concentration factor because of the weld toe undercut. Typical
Large gear that cracked during grinding operations. Localized thermal gradients during grinding resulted in high residual stresses and eventual cracking. Temper etching (dilute nitric acid in water) revealed the presence of abusive grinding.
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Fig. 3
Failure of a mast arm due to fatigue that initiated at a weld toe undercut. Source: Ref 2
causes of undercutting include excessive weld current. The assembly of a group of components can also cause eventual failure. Force-fitting a component creates high residual stresses or damage and causes premature failure to occur. Incorrect placement of a component or incorrect assembly order can also cause high residual stresses or failure to occur. Improper specifications or torque requirements can also cause premature failure. Misalignment of components within the assembly could also result in inadequate service life, because the stresses are not what the designer had anticipated. Service conditions obviously can have a large role in the failure of a component. The service conditions could be normal operations but unanticipated by the designer. It could also be abnormal operations, such as speed, temperature (high or low), or a chemical environment, that were also unanticipated. The lack of proper scheduled maintenance can be a major contributor to premature failure. Maintenance procedures are often reduced as a cost-savings measure. Inadequate lubrication or improper lubrication can also play a role in failure (Fig. 4). In the case of Fig. 4 (Ref 3), the lubrication schedule was extended to reduce aircraft
Fig. 4
The probable cause of this accident was a loss of airplane pitch control resulting from the in-flight failure of the acme nut threads on the horizontal stabilizer trim system jackscrew assembly. The thread failure was caused by excessive wear resulting from insufficient lubrication of the jackscrew assembly. Source: Ref 3
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downtime. This, and other contributing factors, resulted in the loss of 88 lives. Stresses from startup can also contribute, along with rapid temperature gradients and rapid localized changes in the environment. Start-up procedures and maintenance are critical for intermittent operations. Shut-down procedures and resulting stresses are just as critical as proper startup. Inspection procedures to prevent failure are also important. Failure to properly inspect for problems or cracking can be catastrophic (Fig. 5), (Ref 4). In this case, maintenance and inspection personnel failed to detect a fatigue crack in the compressor stage of an aircraft engine. Upon application of power, the compressor stage ruptured, with shrapnel severing fuel lines and igniting the fuel, ultimately leading to the loss of the aircraft.
General Practice Conducting a Failure Analysis The primary objective of any failure analysis is to determine the primary root cause of failure and to establish the appropriate corrective action. There are several stages of an analysis, which can proceed one after the other or occur at the same time. There is no set “fixed-in-stone” procedure, because it is highly dependent on the part and procedures/capabilities of the specific laboratory. These stages of analysis are:
Collection of background data Preliminary visual examination
Nondestructive testing Selection and preservation of specimens Mechanical testing Macroexamination Microexamination Metallographic examination Determination of the fracture mechanism Chemical analysis (bulk and microanalysis) Exemplar testing Analysis and writing the report
These stages are described as follows, and additional information on failure analysis procedures is given in the chapter “General Aspects of Failure Analysis” in this book. Collection of Background Information During the collection of background data, the engineer is trying to gather an understanding of the purpose of the part. The engineer is attempting to discern the design criteria, service conditions, and failure conditions. In the background information, the operating details and manufacturing history should be examined and collected. This manufacturing history should include all the mechanical processing, thermal history or processing, and any chemical process performed on the part. The service history should include all the maintenance records of the part. It should also include the expected environment and loading at the time of failure, as well as the normal environment and loading. Any quality records should be examined for discrepancies. Unfortunately, these records are not always available, and it is often up to the experience of the engineer to determine the quality of the part. Preliminary Visual Examination
Fig. 5
The probable cause of this accident was the failure of maintenance and inspection personnel to perform a proper inspection of a seventh-stage high compressor disk, thus allowing the detectable crack to grow to a length at which the disk ruptured under normal operating conditions, propelling engine fragments into the fuselage. The fragments severed the right engine main fuel line, which resulted in a fire that rapidly engulfed the cabin area. Source: Ref 4
Documenting the failure or fracture is extremely important. There can never be too many drawings or photographs. The cost of photographs (especially digital) is cheap compared to analysis. A high-quality camera with macrocapability is very important and is one of the best tools that a failure analysis laboratory can have. The use of gray cards to ensure proper color rendition is also very important, because the color of scale or oxides can often give an indication of the temperatures that the part has experienced. Sample selection is also very important. All associated debris should be collected and identified. Similar parts should also be collected for
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comparison. In the case of a fastener failure, it is important that the nut and washer be collected, too. All mating pieces should be gathered for subsequent analysis. Any abnormal conditions should be observed and compared with new and used components. Any discoloration or debris should be noted and collected. Any distortion of the part should be noted, along with dimensions of the part. Weather conditions at the time of failure should be collected, as well as all bearing and lubrication conditions and records. During the initial wreckage analysis, the determination of all wreckage should be identified and located on a map or grid before any is touched or moved. Photograph each piece of wreckage and its surroundings. Inventory the parts present or missing. Determine the operating conditions at time of failure. This should include the position of control surfaces, power settings, position of throttles, and any lights or annunciations that occurred. As best as possible during the initial examination of the wreckage, the sequence of failure should be determined. This can be accomplished by examining chevron markings and crack order. The parts should then be closely examined and reassembled. DO NOT allow the fracture surfaces to touch each other, because this can cause potential damage to the delicate surfaces. This analysis can also help determine the sequence of events leading up to failure. Preliminary examination of the part should note any paint, debris, or deposits present. Always remember to “do no harm.” The visual examination should be detailed. Fracture surface crack directions should be noted, identified, and documented. Any abuse or discoloration should be identified, and a general assessment of the workmanship should be determined. Document all findings with photography, with multiple photographs taken from different directions. The incorporation of rulers or scales is important to determine the size and direction of fracture.
Nondestructive Testing Nondestructive testing is very useful for determining the extent of cracking. Magnetic particle inspection is useful for ferrous alloys, with dye-penetrant and ultrasonic inspection as additional methods available for initial inspection.
Magnetic particle inspection uses discontinuities in the magnetic field to identify cracks or discontinuities. Fluorescent dyes with small magnetic particles are used. These magnetic particles gather at the discontinuities in the magnetic field, indicating flaws or indications. It is a common, sensitive, and reliable method that is simple to learn and use. This method has no limitation in part size but is limited to magnetic materials. No elaborate precleaning of the surfaces is necessary. Detection is limited to the surface of the part or section examined. Care must be exercised to prevent local arcing. The dye-penetrant method is useful for examining surface flaws or cracks. It is used primarily for nonferrous alloys but is used for examining ferrous weldments for cracks and porosity. In this method, a high-wetting liquid is spread on the surface of the part. Excess liquid is wiped off. A developer is applied to the part surface. Any cracks, flaws, or other indications will appear. Limitations of this method are the necessity of cleaning the surface prior to and after application of the indication fluid and developer solution. Surface features may also mask indications. It is simple to use, but an understanding of the limitations must be understood prior to application to a part. Eddy-current methods depend on the principle that all metals conduct electricity. An alternating current is applied, and eddy currents occur by electromagnetic induction. Cracks or other flaws cause distortions in the electromagnetic fields, with a result of changing the field impedance. The advantage of this method is that subsurface discontinuities can be detected. No special skill is required to use this method, and the method can be automated. Probe contact with the part is not needed. Limitations of this method are that the depth penetration is limited, and the part must be capable of conducting electricity. Reference standards are needed for specific flaw sizes and materials. Many things can influence readings, including segregation, carburized layers, and changes in profile. Ultrasonic testing uses high-frequency sound waves transmitted through a conducting medium. Any discontinuous boundary can cause a deflection. This method is very sensitive and has high penetration. It is possible to get accurate measurements of flaw position and size, but reference standards must be used. Shape and size can cause errors in interpretation. Experienced operators are required to properly interpret the
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results of testing. Effects of grain size, porosity, and inclusions can also hinder interpretation. Radiography, using x-rays, neutrons, or gamma rays, is also often used to examine structures. Film or sensors (charge-coupled devices) pick up the emitting radiation, with the intensity proportional to the density of the sample. Light areas indicate a dense region, and dark areas indicate a greater exposure or less dense region. Advantages of radiography are the detection of subsurface and internal features at various depths and the documentation of these features by film or other imaging techniques. The primary disadvantage is that reference standards must be used, and the area for testing must be enclosed to prevent radiation from leaking out. Mechanical Testing Mechanical testing is useful to determine the properties of the part and to verify that it meets expected properties and specifications. There are many types of mechanical testing available, including hardness, tensile testing, and impact fracture testing. Hardness testing is probably the most versatile and widely used. It is often used to evaluate heat treatment and can be used as an approximation for tensile strength. It can be used to detect the presence of work hardening or softening and hardening or softening from localized thermal events such as grinding. For the most part, it is a nondestructive test. For microhardness testing, it is necessary to use a metallographic specimen. Tensile testing is used more to establish conformance to specification. It is not necessary to show inadequate ductility because of service loads. Because of the size of the tensile specimen, it may not be possible to excise an appropriately sized sample from the part. Anisotropy of properties can be expected to lower measured tensile and yield strength properties. Impact and fracture toughness testing is typically used to determine conformance to specifications. Charpy impact testing has a high variability in results and may be temperature related. Results must be taken with temperature in mind and may not correlate with real results because of size limitations. Fracture toughness testing and the results from KIc testing can be used in design, and the results are useful for calculating critical flaw sizes. It can also be used to examine estimated crack growth rates; however, samples are difficult to prepare and
test. These methods also do not incorporate the effects of residual stresses. Selection and Preservation of Specimens The selection and preservation of fracture surfaces is vital to prevent the destruction of evidence. Unprotected, the fracture surfaces or parts can become mechanically or chemically damaged. This damage can obliterate evidence and make the determination of fracture difficult or impossible. Both sides of the fracture must be protected. This is in the event that if one surface is damaged, the other side can be examined. Protection of the specimens during shipment is also very important, because evidence could be destroyed. Avoid touching surfaces with the hands, because the chemicals and acids present can cause artifacts or destroy data. NEVER fit surfaces together, because the delicate fracture features can be destroyed. Since both surfaces would be damaged, it could destroy the chances for determining the fracture mechanism. Cleaning of specimens is to be done only when absolutely necessary. For the most part, it is required to prepare the sample for the scanning electron microscope (SEM). Dry air blasts or soft artist brushes are typically all that is needed. Rinsing in organic solvents then evaporating the solvent with dry air is useful for preparing specimens for the SEM. Chemical cleaning is generally not recommended under any circumstance. Foreign substances such as scale or debris should be preserved. Do not use rust inhibitors, because of the inevitable damage to the part and fracture surfaces. These rust inhibitors are also extremely difficult to remove. Avoid washing the sample or parts with water unless seawater or other chemical is present. In this case, gently wash with distilled water and follow that with high-quality alcohol or acetone. Allow to dry and place in a dessicator. Plastic replicas are useful in preserving fracture surfaces and removing debris for further analysis. Softening replica tape (available at transmission electron microscope supply houses) with small amounts of acetone forms plastic replicas. The softened tape is pressed gently onto the fracture surface. Additional layers of tape, softened with acetone, are applied to the fracture surface. After multiple layers have been applied, the entire replica is allowed to dry and then is placed in a dessicator. When the part is ready to be examined, the replica is
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carefully removed using tweezers. Any debris on the surface is also preserved for further analysis in the replica. Multiple plastic replicas can be used to clean a surface of a part. This can be repeated as necessary. Sectioning of Specimens Sectioning is very important, because it captures the portion of the fracture surface for examination or the appropriate metallographic specimen. The biggest limitation is size. It is important that the portion to be removed is documented by photographs and sketches, showing the location of the specimen to be removed. Preserve any fracture surface by plastic replicas or other method to prevent damage or attack. Regions adjacent to cracks are also to be preserved and protected. Cutting the specimens should be done very carefully so as not to cause any heat damage. Coolants are not recommended, unless the material cannot be cut without heat generation. The use of plastic replicas is useful for protecting surfaces and preserving any debris present. Opening secondary cracks is useful when the primary fracture surface is damaged. These secondary cracks may provide better information, because they are tightly closed, and the fracture surfaces are not exposed to surface contaminants and corrosion. Care must be taken not to damage the primary fracture surface. Bending to open the crack is preferable, to expose the crack face. Often, the use of a sawcut to the back of the part will reduce the amount of force necessary to open the secondary crack. Another method is to use a tensile machine to open the crack face. The crack opening should be measured prior to opening, and the crack opening displacement can also be measured as the crack is slowly opened and exposed. One technique is to immerse the specimen in liquid nitrogen and impact the part so that the fracture surfaces are rapidly opened. One problem with this method is that it is very easy to damage the fracture surface from a misapplied hammer hit.
information on the location of fracture origins, direction of cracking, configuration of the stress state, and the last region to fail (shear lip). The presence of chevron marks can indicate the direction of rapid crack growth, and the different textures of the fracture can differentiate between fast final fracture and the initiating mechanism of fracture. Different textures from the region of fast fracture can indicate a different mechanism, such as fatigue, stress-corrosion cracking, or hydrogen embrittlement. Microscopic Examination The microscopic examination is usually conducted with an SEM (Fig. 6). This instrument is probably the most useful of all instruments for determining the mechanism of failure. It is capable of a large depth of field, with magnifications of 10 to 300,000 · . It allows for direct examination of specimens, and when coupled with an energy-dispersive spectrometer, very small regions can be examined and analyzed for chemistry. It is very easy to use and requires very little training to take quality images. Interpretation of the images requires experience and understanding of the four basic modes of failure: dimpled rupture, cleavage, brittle intergranular, and fatigue. From these four basic modes, the detailed mode can be examined, and the failure mechanism is fit to the evidence. A greater discussion of the mechanisms of failure is found later in this chapter and elsewhere in this book. Metallography Metallography is a vital part of a failure analysis investigation. It can examine crack
Macroscopic Examination The macroscopic examination is conducted by a detailed examination at 1 to 100 · by eye or binocular microscope. High-quality optics with excellent depth of field are required to properly examine the fracture surfaces. This detailed macroscopic examination can reveal a wealth of
Fig. 6
Typical scanning electron microscope used for microscopic analysis of a fracture surface
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morphology and its relationship with the microstructure present. It can help determine the thermal history of a component or region of a part and can show if work hardening was present. There can never be too many photographs and metallographic sections. Metallographic sections should be taken away from the crack and near the determined origins of cracking. Because this method is destructive, it is undertaken last. Typically, the crack face and edges are protected from rounding by applying support. This support can be electroless nickel plate or the use of alumina beads or steel shot in the metallographic specimen, adjacent to the surface. Metallographic specimens are prepared using an epoxy or phenolic resin. The sample is placed into a small press, and phenolic resin is poured over the section. The press compacts the resin and forms a small, round sample that is then polished, etched, and examined under a metallographic microscope. When the specimen has cooled, it is taken out of the press and ground through a sequence of sandpapers. Typically, the sequence is 240, 320, 400, and 600 grit. The specimen is ground very flat before polishing. During polishing, the metallographic specimen is polished using a flat platen and 3 mm alumina slurries. Final polish is accomplished using 0.15 mm alumina slurry. Other polishing agents can be used, with diamond being a very common polishing agent. A finished metallographic sample used for the determination of the fracture mechanism in a steel weldment is shown in Fig. 7. Examination of the metallographic specimen reveals surface imperfections, inclusions, and microstructural details. It can reveal the presence of decarburization and improper heat
treatment. It often provides the needed documentation and support for the fracture analysis and determination of the root cause of failure.
Determination of the Fracture Mechanism Examination of the fracture surface and metallography are used to determine the cause of failure. First, it is necessary to determine the fracture mode. Unfortunately, there is no clear or logical classification of fracture. Generally, classification is based on the crack growth mechanism (see also the chapter “General Aspects of Failure Analysis” in this book). Ductile Fracture On a macroscopic scale, a ductile fracture is accompanied by a relatively large amount of plastic deformation before the part fails. After failure, the cross section is reduced or distorted. Shear lips are observed at the latter part of the fracture and indicate the final failure of the part. The fracture surface is dull, with a fibrous appearance. Microscopically, ductile fracture is characterized by several distinct stages (Ref 5–8); an example is shown in Fig. 8. In this case, an ISO 12.9 low-alloy bolt failed by ductile torsional overload. The fracture was smooth, with fracture initiating from the threads. The fracture mode was microvoid coalescence (Ref 9), which occurs by the following process:
Fig. 7
Typical metallographic specimen. This specimen was used to examine microstructures in a failed weldment.
A free surface is created from a small particle. This particle can be a second-phase particle, dispersoid, or inclusion. The separation of the metal matrix from the small particle at the matrix/particle interface can form this free surface, or the fracturing of the small particle can form the free surface. The free surface around the small particle creates a void. This void grows by plastic strain and hydrostatic stress. Finally, the voids grow to a size that they join or coalescence with adjacent voids.
This process of void formation, growth, and coalescence is shown schematically in Fig. 9. If the particles are well matched to the matrix and form a strong interface between the matrix and the particle, then the initial formation of voids is the critical step. Fracture occurs shortly after
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Fig. 8
Fracture of an ISO 12.9 bolt by ductile torsional overload. (a) Overall view of fracture. (b) Smooth and fibrous fracture as seen through the SEM. (c) Microvoid coalescence (dimples)
Fig. 9
Schematic showing the formation of microvoid coalescence
void formation (Ref 10). If the interface between the particles and the matrix is weak, then voids form and grow readily. Substantial plastic deformation occurs. Fracture occurs when the voids reach a critical size. These voids substantially reduce the cross section, with the resulting local plastic instability (Ref 11). These voids coalesce to form a central crack perpendicular to the applied tensile stress. Depending on the applied stresses, the shape and configuration of the dimple shape can be changed (Fig. 10). This fact is important in determining the type of loading during a postfracture investigation. Dimples are small and can only be detected by using electron microscopy (Fig. 11). The presence of inclusions in steel plays a major role in the ductility of steel. As indicated previously, the inclusions fracture and separate from the matrix during decohesion. Therefore, the deformability of these inclusions is important to determine the ductility of steel. Nearly all steels have nonmetallic inclusions. The size and frequency of these inclusions is determined by the methods described in ASTM E45 (Ref 12). The cleanliness of the steel is
Fig. 10
Schematic representation of the creation of dimples in a loaded member by (a) simple tension, (b) shear loading, and (c) tearing
Fig. 11
Microvoid coalescence as seen through the SEM
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important to the ductility of the steel. All other things being equal, the steel with the lower inclusion size, shape, and frequency will have a greater ductility than another steel with a greater inclusion count. Modern steelmaking practices generally produce low inclusion content. Often, steels for aerospace applications require a frequency/severity determination of inclusions in accordance with AMS 2300, AMS 2301, AMS 2303, or AMS 2304 (Ref 13–16). A specificsized test specimen must be heat treated and examined using magnetic particle inspection. The procedures are outlined in the aforementioned specifications. The inclusions found in steels have been divided into five categories related to their deformation behavior (Ref 17):
The inclusions Al2O3 and calcium aluminates are produced during deoxidation of steel during the production of molten steel. They are brittle at practically all temperatures. Spinel-type oxides are not deformable up to 1200 C but may be deformed above this temperature. Silicates of calcium, manganese, iron, and aluminum in various proportions are brittle inclusions at room temperature but become more deformable at higher temperatures. The formability increases as the melting temperature of the silicate decreases. Therefore, aluminum silicate has much less formability than the lower-melting manganese silicates. FeO and (FeMn)O are deformable at room temperature but gradually become more brittle at temperatures above 400 C. Manganese sulfide (MnS) is the most common inclusion found in steel, and it is increasingly deformable as the temperature falls. The morphology of the MnS inclusions changes, depending on how they were formed.
Ductile failure can occur with any of the types of inclusions. This is true whether it is the brittle alumina-type inclusions or the more ductile sulfide-type inclusions. Inclusions generally initiate ductile cracking above a critical size. Coarser inclusion sizes tend to have a larger local stress-concentration factor, which can cause local decohesion and microcrack formation. Work by Maropoulos and Ridley (Ref 18) has shown the effect of volume fraction of ironalumina on the ductility of steel. Increasing
amounts of inclusions reduce the ductility of the steel. A reduction in the yield stress, due to the stress concentrations around the inclusions, is evident at low volume concentrations of inclusions. The presence of inclusions in the size range of 1 to 30 mm reduces the energy absorbed during ductile fracture. Fine dispersions of ductile inclusions will delay the onset of cleavage-type fracture by localized relaxation of stresses. At the same time, the yield stress also increases. During deformation, forming, or forging, the ductile inclusion MnS has a marked effect on the ductility of the final product. Types 1 and 2 MnS inclusions will elongate on deformation, while type 3 MnS inclusions will rotate into the rolling plane. This will reduce toughness and ductility in the transverse direction. Type 2 inclusions are the most harmful to ductility and toughness, so some effort is being made to eliminate these inclusions by ladle additions of other strong sulfide formers, such as titanium, zirconium, and calcium. Ductility is also influenced by the fact that MnS contracts more than the iron matrix upon cooling. The bond between the MnS inclusion and the matrix is not strong enough to prevent microvoid formation. Because MnS inclusions tend to form as strings or stringers along the rolling direction, the toughness and ductility are strongly influenced in the rolling direction. Transverse to the rolling direction, ductility and toughness are much worse. In a similar fashion to that of inclusions, the distribution of carbides can also influence the toughness and ductility of the steel. The strain needed for void formation decreases with increasing carbide volume fraction. Spheroidal carbides will not crack at small strains and exhibit decohesion. Spheroidized steel is much more ductile than similar steel of the same hardness containing only ferrite and pearlite. Pearlite has a lower critical strain for void formation. In addition, when a crack or void forms in a pearlitic matrix, it will tend to run along the length of a pearlite lamella. Examining this type of fracture under the SEM reveals that the base of the dimples contain fractured pearlite lamella. Brittle Fracture Very little plastic deformation and a shiny fracture surface characterize brittle fractures. Often, chevron patterns point back to the origin
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of failure (Fig. 12) (Ref 19). It can occur at low stress and propagate with rapidity, often at speeds approaching the speed of sound in the failed material. Since the early 1940s, there has been tremendous growth in the number of large welded structures. Many of these structures have failed catastrophically in service, most notably the “Liberty ships” (Ref 20) used to transport war material during World War II. Analysis of the fracture surfaces of the failures (Ref 21) indicated that they initiated at a notch and propagated with no plastic deformation. These notches were of three types: Design features: Structural members were rigidly joined at angles less than 90 and then welded. Fabrication details: Procedures used during the manufacture of the part caused the formation of notches. Welding arc strikes, gouges, and fitting procedures created physical notches. Welding procedures and
In brittle fractures, limited energy is absorbed by the fracture. Energy is absorbed through regions of small plastic deformation. Individual grains separate by cleavage along specific crystallographic planes. This is shown in Fig. 13. Visually, little or no plastic deformation or distortion of the shape of the part characterizes brittle fractures. The fracture is usually flat and perpendicular to the stress axis. The fracture surface is shiny, with a grainy appearance. Failure occurs rapidly, often with a loud report. Because the brittle cleavage is crystallographic in nature, the fracture appearance is faceted. Often, other features are present, such as river patterns (Ref 23). These are shown schematically in Fig. 14. There are three basic factors that contribute to a cleavage type of fracture in steels. They are:
Fig. 12
Fig. 13
Chevron markings point back to the origin of failure in brittle steels. Source: Ref 19
heat treatment caused metallurgical or microstructural notches to occur from abrupt changes in microstructure or the production of microstructures that were brittle. Features such as porosity from welding or casting also caused brittle fracture initiation. Material flaws: These flaws resulted from melt practice at the mill and appeared as large inclusions, internal oxidation, porosity, or segregation.
Triaxial stress state that forms at a notch, similar to that described previously Low temperature High strain rate or rapid loading rate
These three factors do not have to be present for cleavage-type fracture to occur. Most brittle, cleavage-type fractures occur when there is a triaxial stress state and low temperature. This is
Cleavage fracture in a low-carbon steel, seen through an SEM. Cleavage fracture in a notched impact specimen of hot-rolled 1040 steel broken at 196 C ( 320 F), shown at three magnifications. The specimen was tilted at an angle of 40 to the electron beam. The cleavage planes followed by the crack show various alignments, as influenced by the orientations of the individual grains. Grain A, at center in fractograph (a), shows two sets of tongues (see arrowheads in fractograph b) as a result of local cleavage along the {112} planes of microtwins created by plastic deformation at the tip of the main crack on {100} planes. Grain B and many other facets show the cleavage steps of river patterns. The junctions of the steps point in the direction of crack propagation from grain A through grain B, at approximately 22 to the horizontal plane. The details of these forks are clear in fractograph (c). Source: Ref 22
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actuated by a high rate of loading. Many types of tests have been developed to determine the susceptibility of steels to brittle behavior. These tests include the Charpy impact test (ASTM E23) (Ref 24) and the fracture toughness test (ASTM E399) (Ref 25). Others include the nilductility test (ASTM E208) (Ref 26) and dynamic tear test (ASTM E604) (Ref 27). The notch toughness of low- and mediumstrength steels is highly dependent on temperature. There is a transition from ductile fracture to brittle fracture as the temperature decreases. One criterion for the transition temperature is the nil-ductility temperature (NDT). The NDT is the temperature where fracture becomes 100% cleavage, and there is essentially no plastic deformation.
Fig. 14
Changes in the NDT can be produced by changes in microstructure and chemistry. The largest change can be effected by changes in the amount of carbon and manganese. The NDT is lowered by approximately 6 C (10 F) for every 0.1% increase in the manganese concentration. Increasing the carbon content also lowers the NDT. The manganese-carbon ratio should be approximately 3 to 1 for good notch toughness. Decreasing the concentration of phosphorus also decreases the NDT. Nitrogen causes the NDT to increase (more brittle). However, because of the interaction with other alloying elements in steel, it is difficult to quantify the increase of NDT with increasing nitrogen concentration.
Schematic of river patterns formed in brittle materials. (a) Tilt boundary. (b) Twist boundary. Source: Ref 23
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Nickel is beneficial for increasing ductility. Up to 2% Ni is effective in lowering the NDT. Increasing concentrations of silicon have the effect of increasing the NDT. Chromium has nearly no effect, while molybdenum is extremely effective in increasing the ductility of steels and drastically decreasing the NDT. Oxygen strongly decreases the ductility. It can also cause an increased propensity for intergranular fracture by creating brittle oxides at the grain boundaries. Decreasing the grain size has a strong effect on increasing the ductility and notch toughness. Section thickness can also influence ductile and brittle behavior (Ref 28). The results showed that there was considerable variation of toughness with the thickness of the specimen (Ref 29, 30). Further, at large thickness, the toughness appeared to reach a constant value (Fig. 15) (Ref 31). Within this curve, there are three apparent regions. First, there is the region where maximum toughness is obtained (thin sections). Second, there is the region of intermediate toughness, and lastly, a region with relatively constant toughness (thick sections). In the first region, the fracture appears to consist entirely of a shear lip, or, in other words, the fracture surface is inclined at an angle of approximately 45 to the tensile axis. In this situation, the stress in the direction of the thickness of the specimen tends toward zero, and a state of plane stress is achieved. As the specimen is pulled, it experiences buckling. Because of this buckling, yielding occurs on the through-thickness planes at an angle of 45 to the tensile axis. Crack extension occurs by sliding. This sliding motion is achieved by the movement of a number of screw dislocations (Ref 32, 33) on the 45 plane, as shown in Fig. 16.
In the intermediate range, the fracture behavior is complicated. The fracture does not consist of entirely slant-type fracture, nor does it contain entirely a flat plane-strain-type fracture. Instead, the regions of flat and slant fracture are approximately equal. At the thin end of the thickness range, the slant ligaments on either side of the testpiece carry most of the load. At the thick end of the range, the side ligaments carry a much smaller percentage of the load. The amount of flat fracture increases. This is shown schematically in Fig. 17. It has been found (Ref 28) that the amount of flat fracture depends only on the thickness of the test specimen and was independent of crack length. In the third region, the fracture consists of predominantly flat fracture. Some evidence of very small shear lips may be present at the later part of fracture. Fracture is catastrophic and rapid. No plastic deformation is evident. In this third region, any increase in the thickness of the testpiece causes no further decrease in the toughness. These fracture patterns are useful in determining the state of stress within a failed component and can help to understand the mechanism of failure. One famous failure involving brittle fracture was the “Great Boston Molasses Disaster”
Fig. 15
Fig. 16
Variation of toughness with thickness
Mode of separation in a thin sheet
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(Ref 34). In this failure, the United States Alcohol Company fabricated a large cast iron molasses tank in Boston in December 1915. This tank was 27 m (90 ft) wide and 17.7 m (58 ft) tall, with a head of 15 m (49.5 ft) of molasses. It was fabricated of cast iron plates riveted together. It held 8.7 · 106 L (2.3 million gal) of molasses, ostensibly used for the fermentation of ethanol used for liquor. The man who oversaw construction could not read blueprints, nor did
he have any technical training. No engineers or architects were consulted to ensure that the tank was constructed safely. On January 15, 1919, the tank exploded with great force, and the streets of Boston were flooded with waves of molasses from 2 to over 4 m (8 to 15 ft) tall (Fig. 18). This great wall of molasses was reported to have moved at speeds up to 35 miles (56 km) per hour and devastated a large section of Boston along Commercial Street between Copps Hill and the playground of North End Park. Half-inch steel plates were torn apart, and these plates were thrown with enough force to cut girders of the elevated railway. This explosion, and the subsequent wave of molasses, resulted in 21 people killed, 150 people injured, many buildings destroyed, and an entire area devastated. The elevated train trestles were knocked over. Early accounts of the disaster included reports that the tank was destroyed by anarchists. In a trial, it was found that the company was liable for $628,000 in damages (in 2007 dollars, approximately $7,000,000). Investigation many years later indicated that the probable cause was brittle fracture of the tank at the rivets, with the temperature below the ductile-to-brittle transition temperature. One interesting result of this disaster was that Massachusetts and many other states created laws to certify engineers and to regulate construction. It also required stamped drawings certifying that an engineer had reviewed the plans. It was this failure that was the origin of the professional engineer’s license and stamp, as it is known today (2007). As a side note, the 18th Amendment was ratified and Prohibition signed into law on January 16, 1919. In another example of brittle fracture, an AISI 4330V hook-point, used for the arrestment of
Fig. 17
Schematic of fracture in the intermediate range
Fig. 18
The Great Boston Molasses Disaster. Twenty-one people were killed and over 150 buildings destroyed as the result of 2.3 million gal of molasses flooding North Boston.
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Fig. 19
Arresting gear hook-point, manufactured from AISI 4330V, that failed during landing. Failure occurred at the inner fillet radius of the right-hand lug
naval aircraft on landing, failed during field trials during the 13th arrestment. The landing configuration was severe, with high aircraft sink rates, high aircraft gross weight, and landing at a large angle to the cable. The hook-point failed at the inner fillet radius of the right-hand lug (Fig. 19). The hook-point successfully engaged the arrestment cable, with no other aircraft damage. The part was forged, machined, heat treated, and hard surfaced in the cable groove, using a high-velocity oxyfuel coating for wear resistance. Examination showed that the microstructure of the hook-point was quenched and tempered martensite. Hardness measurements showed that the hook-point had a substantially higher hardness (HRC 54) than the specified hardness of HRC 46 to 48. The chemistry of the hook-point indicated that it was at the high side of the specification, increasing the hardenability of the steel and increasing the resistance to tempering. Hydrogen measurements indicated that the hydrogen content was 0.2 ppm. The high strain rate during landing and the low concentration of hydrogen precluded failure by hydrogen embrittlement. An SEM examination of the fracture surface showed that the fracture contained microvoid coalescence and quasicleavage, suggestive of brittle failure (Fig. 20). Charpy impact testing showed that the impact toughness of the as-received part was significantly lower than a part of the same chemistry properly tempered to HRC 46. Finite element analysis showed a high localized stress concentration at the lug inside fillet radius. It also showed that the stresses were highly triaxial. Based on the analysis, it was determined that the hook-point lug failed by quasi-cleavage,
Fig. 20
SEM fractographs showing (a) location of origin at the inner fillet radius and (b) quasi-cleavage evident on the fracture surface
and that the failure was aggravated by high local stress concentration at the fillet radius, improper heat treatment (making the material more
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brittle), and extremely high dynamic loading. It was recommended that the radius be made larger to reduce the stress concentration and also to retemper the hook-points to meet specification. Intergranular Brittle Fracture Another form of brittle fracture is called intergranular cracking. In this fracture mechanism, failure occurs by decohesion along grain boundaries and not on specific crystallographic planes, such as in cleavage fracture. Intergranular cracking can have several different causes. Typical causes of intergranular cracking in steel alloys include:
Quench-age embrittlement: Cooling of carbon steels and low-alloy steels from subcritical temperatures can precipitate carbides within the microstructure. The strength is raised, but toughness is lost. Quench cracking: During quenching, the transformational and residual stresses developed during quenching of steels can cause cracking during heat treatment. Tempered martensite embrittlement: Within the range where blue-purple oxides can form on steels (230 to 370 C, or 450 to 700 F), precipitates can form that increase the tensile strength and hardness while reducing the ductility and toughness. Temper embrittlement: Quenched steels containing appreciable amounts of manganese, silicon, nickel, or chromium are susceptible to temper embrittlement if they contain even trace amounts of antimony, tin, or arsenic. Embrittlement of susceptible steels can occur after heating in the range of 370 to 575 C (700 to 1070 F) but occurs most rapidly at approximately 450 to 475 C (840 to 885 F). Graphitization: This happens when the pearlite in steels begins to decompose into ferrite and graphite following very long, high-temperature service, for example, in steam power stations. For these applications, a few steels turn out to be satisfactory, while many others are subject to graphitization. Internal oxidation: This is one of the common failures in high-temperature, oxidizing conditions. Liquid metal embrittlement or solid metal embrittlement: Intermetallic compounds form at grain boundaries when low-meltingtemperature metals (cadmium, zinc, etc.)
penetrate by diffusion. An example would be galvanized steel where the zinc has diffused into the steel in the vicinity of 420 C (787 F). Hydrogen embrittlement: The presence of hydrogen and static loads or a low strain rate can result in hydrogen embrittlement. Stress-corrosion cracking Grain-boundary decohesion at elevated temperatures (creep rupture)
The fracture surface appearance of intergranular cracking is generally shiny and faceted. It has a “rock-candy” appearance. Often, when the mechanism is from corrosion, the corrosion product is present. This can dull the appearance of the facets. The appearance of intergranular fracture is most clearly seen in the electron microscope, and an example is shown in Fig. 21. Quench cracking is the limiting case of excessive residual stresses exceeding the tensile strength of the material. Two processes contribute to quench cracking, as well as distortion and residual stresses. The first process is the stress from the volume expansion of martensite during transformation from austenite to martensite. The second source is from thermal stress due to differential contraction due to different cooling rates in the steel. The transformational stress from the formation of martensite is primarily responsible for cracking during quenching, and thermal stresses from differential cooling are usually from subcritical heat treatments such as annealing. During quenching, the volume expands from the close-packed face-centered cubic structure of austenite to the body-centered tetragonal structure of martensite. This volume expansion
Fig. 21
Intergranular fracture from hydrogen embrittlement, as seen through the SEM
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is approximately 4% and is related to the carbon content of the steel. During quenching, the outer surface of the part cools first and transforms to martensite. There is an attendant volume expansion at the surface, and the untransformed and still hot interior surface usually has sufficient plasticity to accommodate the changes in the part volume. The outside surface is in compression. Upon cooling, the interior of the part also transforms to martensite but is constrained by the hard outside surface layer of previously transformed martensite. On the transformation of the inner core, a volume expansion occurs in the interior of the part, and the outer surface is placed in tension. If quenching is severe, the resulting tensile residual stresses can exceed the ultimate tensile stress of the surface untempered martensite. Cracking is intergranular and often exhibits an oxide scale on the fracture surface. If cracking occurred during quenching, remnants of quench oil can be found on the surface of the crack, and often, elevated-temperature scale is apparent. Cracking can be delayed due to the transformation of retained austenite. This is one reason why it is recommended to temper parts immediately after quenching. Should delayed quench cracking occur, then the temper scale is thinner and often shows the characteristic temper colors, indicative of the temper temperature. High-carbon steels and steels with high hardenability are the most prone to quench cracking. Surface features such as sharp radii, large changes in section, or the presence of laps, burrs, rough-machined surfaces, and other surface discontinuities increase the constraint during quenching and increase the propensity toward quench cracking. Quench cracking can be mitigated by improved surface condition and the removal of scale, burrs, and sharp edges. Geometry changes, by increasing transitions from thin to thick sections, and generous radii can also help reduce quench cracking. The use of higherhardenability alloys will also reduce the propensity for cracking, because it will allow a reduced quench rate to achieve the same properties. Reducing the austenitizing temperature or reducing the temperature differential between the austenitizing temperature and the quenchant temperature will reduce the propensity for cracking. Often, the geometry is set, as is the alloy of the part. In this case, the heat treater can reduce the quench rate or use martempering to reduce quench cracking.
Martempering is the process of using hightemperature quench oils and quench oil temperatures of 90 to approximately 200 C (200 to 400 F). The part is quenched into the hightemperature oil, and the parts are allowed to equilibrate or at least minimize the temperature gradient across the interior of the part. The part is then removed from the oil and allowed to cool in any convenient manner. This method has proven to be very effective in reducing quench cracking as well as distortion from quenching. A long pinion gear failed in service near the midlength of the shaft (Fig. 22). One gear tooth fractured during service, resulting in the gear being removed from service and sent to the laboratory for failure analysis. Magnetic particle inspection, using a fluorescent dye, revealed the presence of multiple linear indications on cracking of the gear tooth faces (Fig. 23). Examination of the fracture surface showed a discolored region at the origin of cracking (Fig. 24). This discolored region was attributed
Fig. 22
As-received pinion gear that failed in service
Fig. 23
Magnetic particle inspection of the failed pinion gear showed arc-shaped cracks on the gear tooth faces.
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Region of Cracking
Fig. 24
Overall view of the cracked pinion showing the location of the fracture and the presence of a discolored region
Smeared surface showing secondary cracking Region of Discoloration
Fracture Origins
Region of Discoloration
Fig. 26
Rough machining at the surface of the tooth showing smearing and tearing of the machined surface. This is suggestive of abusive machining, due to dull cutting tools, inadequate coolant, or excessive speeds and feeds.
Fig. 25
Closeup of the fracture region showing the discolored region. The color of the oxidation indicated that the crack occurred after quenching and during the tempering operation.
to oxidation that occurred during heat treatment. The coloration of the oxide scale suggested that the oxidation occurred during
tempering (Fig. 25, 26). If the crack was preexisting prior to heat treatment, it would be darker and thicker. Examination of the tooth faces showed secondary cracking at regions of tearing and smearing along the tooth face (Fig. 27), suggestive of abusive machining practice, including the use of a tool that was dull or excessive feeds and cutting speeds.
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Fig. 27
Secondary cracking evident at regions of abusive machining
Fig. 28
Metallographic specimen of the pinion showing inadequate case at the root of the tooth. Etched in
0.5% nital
Fig. 29
Large gear that showed evidence of cracking. (a) As-received gear. (b) Crack evident on gear face. (c) Region after temper etching showing evidence of abusive grinding
Metallography of the teeth showed no evidence of burning or excessive temperature. The root of the tooth showed little evidence of proper hardening or case (Fig. 28). The microstructure in the root consisted of ferrite and pearlite, with lightly tempered martensite, further suggesting inadequate heat treatment. The tooth tip showed a fine-grained martensitic structure. No evidence of overheating was present. Examination of the tooth surface showed tears and smearing. Microhardness of the hardened regions of the tooth showed a hardness of HRC 58, while the root of the tooth was HRC 29, consistent with the observed microstructure. Investigation of the induction heat treating conditions revealed that the concentration of the quenchant used was approximately 5%, while 6 to 10% was specified. The concentration was controlled solely by refractometer. Contamination of the quenchant was unknown. Based on the evidence, it was determined that fracture and failure of the pinion gear tooth was
caused by quench cracking, aggravated by improper concentration control and inductionhardening parameters. The situation was further aggravated by poor machining practice, creating tearing and smearing at the surface. Often, quench cracking can result not from heat treating operations but from other sources, such as abusive grinding (Fig. 29). In this case, a large gear was found to be cracked. As is usually the case, the heat treater was blamed. Temper etching of the region of cracking showed a darkened region, suggesting overtempering of the part in a localized region. Localized overheating during service can also result in quench cracking. A hook-point, used for catching the large cable on an aircraft carrier, showed evidence of cracking in the cable groove (Fig. 30) after a carrier landing. The hook-point was manufactured from AMS 6411 (AISI 4330V), heat treated, and a high-velocity oxyfuel (HVOF) coating was applied. Imprints of the arresting cable were left in the cable groove.
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The vertical cracks were exposed using liquid nitrogen and impact loading. The crack faces were discolored with a golden tint; they were subsequently examined using the SEM (Fig. 31). The cracks were intergranular along
Fig. 30
Arresting gear hook-point showing vertical cracking in the cable groove and evidence of localized heating (from the temper colors in the cable groove)
Fig. 31
prior-austenite grains. The laboratory fracture showed microvoid coalescence. The cracks showed three distinct regions: incipient melting at the surface, intergranular regions, and laboratory-induced ductile fracture (Fig. 32). A metallographic section (Fig. 33) through the vertical cracks showed untempered martensite at the surface, a transition region of overtempered martensite, and finally, a region of tempered martensite. No evidence of the HVOF coating was observed at the crack initiation site. Hardness in the core was KHN 460. Hardness in the transition region was KHN 390, and the surface had a hardness of KHN 620. The microstructure is similar to a weldment heataffected zone and shows that a significant heat event occurred. Chemical analysis showed that the material conformed to AMS 6411, with alloying elements at the top of the range increasing the sensitivity to quench cracking. Hydrogen analysis indicated 0.8 ppm hydrogen. The levels of hydrogen present and the high strain rate of
SEM examination of the vertical cracking. “A” indicates the presence of intergranular cracking along prior-austenite grain boundaries. “B” indicates microvoid coalescence from the laboratory fracture during the exposure of the crack face.
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loading precludes hydrogen embrittlement as a possible failure mechanism. Based on the analysis, it was determined that the vertical cracking in the cable groove was the
Fig. 32
Fig. 33
result of transformational stresses from frictional heating during capture of the arresting cable by the hook-point. The mechanism is similar to quench cracking and was aggravated
Exposed crack face showing two distinct regions on the crack face. “A,” region of incipient melting. “B,” intergranular fracture
Metallographic section through the vertical crack showing (from right to left) a lightly etching region of fine-grained untempered martensite, a transition region of overtempered martensite, and a region of nominally tempered martensite. Hardness in the untempered martensite was KHN 620. The transition region showed a hardness of 390 KHN, and the nominal core hardness was KHN 460
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by the higher-than-normal hardenability of the alloy. Recommendations included the application of a different HVOF coating to better resist the frictional heating of the cable during carrier arrestments. Tempered martensite embrittlement (TME) may not be associated with impurity atoms segregating to prior-austenite grain boundaries. The most common factor in TME is the formation of cementite during tempering (Ref 35). When a given steel has a low impurity content, the source of TME is the decomposition of retained austenite during the second stage of tempering. Thomas first proposed this mechanism (Ref 36). This was found when transmission electron microscopy showed the presence of thin regions of retained austenite between martensite in as-quenched steels, which subsequently transformed to cementite on tempering in the range of 230 to 470 C (450 to 700 F). The presence of phosphorus also plays a role. If two steels are compared, one containing a higher concentration of phosphorus, the steel with the higher phosphorus content will have poorer impact properties than an identical steel with a lower phosphorus level. This will remain true through the entire range of tempering temperature up to approximately 500 C (932 F). The fracture mode is intergranular along prioraustenite grain boundaries (Ref 37). It is likely that phosphorus is present at the prior-austenite grain boundaries. It is only after cementite precipitates in the tempered martensite that TME is fully present. Often, the presence of molybdenum at concentrations up to approximately 0.5% will reduce the effect of TME. On June 19, 1974, during a cold start after a long shutdown for repairing the Tennessee Valley Authority Gallatin No. 2 unit, the intermediate-pressure/low-pressure rotor burst at approximately 3400 rpm. The rotor had been in operation for 106,000 h from its operational start in May 1957 (Ref 38). The burst rotor was forged from an air-melted ingot. This ingot was produced by a large region of MnS segregation zone that was present at the center of the ingot, which was subsequently bored by machining during fabrication of the rotor. The steam temperature was 566 C (1050 F). Tempered martensite embrittlement occurred over the long period of operation and substantially reduced the toughness of the rotor. The presence of the MnS inclusions initiated fracture by creep-fatigue interaction and was enhanced by the presence of TME (Ref 39).
Rail steels have been documented to fail because of TME (Ref 40). This was especially true of older rails manufactured in open-hearth furnaces with high phosphorus content. This occurred because of slow cooling through the 500 C (930 F) range or from isothermal holding at 500 C. Figure 34 shows a representative SEM fractograph of an Fe-0.26C-2.11Si2.27Mn-1.59Cr wt% carbide-free bainitic rail steel that has been temper embrittled by heat treatment at 500 C for 5 h (Ref 40). Temper embrittlement is only now becoming understood with regard to its mechanism. However, the conditions of temper embrittlement are well known (Ref 41, 42). Steels must be heat treated or cooled through the range of 375 to 575 C (706 to 1070 F) in order to become temper embrittled. Temper embrittlement is typically detected by an increase in the ductile-to-brittle transition temperature. This is shown in Fig. 35 for AISI 3140 steel temper embrittled by furnace cooling through the critical range and holding at 550 C (1020 F) (Ref 35). The embrittlement reaction follows a typical C-curve, with the minimum in embrittling time at approximately 1 h at 550 C (1020 F) and several hundred hours at 375 C (706 F) (Ref 43). By heating to approximately 575 C (1070 F), temper embrittlement is reversible and can be eliminated after holding for only a few minutes at temperature. For temper embrittlement to occur, specific embrittling impurities must be present. These include antimony, phosphorus, tin, and arsenic. Quantities of less than 0.01% are enough to cause temper embrittlement. For the most part, simple plain carbon steels are not considered to be susceptible to temper embrittlement
Fig. 34
SEM fractograph of Fe-0.26C-2.11Si-2.27Mn1.59Cr wt% carbide-free bainitic rail steel that has been temper embrittled by heat treatment at 500 C for 5 h
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as long as manganese concentrations are held to below 0.5%. Alloy steels containing chromium and nickel are the most prone; however, additions of molybdenum at a concentration of up to 0.5% are effective in reducing the susceptibility of these steels to temper embrittlement. Large forgings have been prone to temper embrittlement because of the slow cooling that occurs during fabrication. These large forgings are also prone because of the operating temperatures applied, especially in large turbine rotors. Liquid Metal Embrittlement or Solid Metal Embrittlement. Exposure of steels to liquid metals has been observed to result in brittle fracture along prior-austenite grain boundaries (Ref 44). Steels may be embrittled by exposure to any of the low-melting metals shown in Table 1 (Ref 45). Embrittlement occurs by wetting of the prior-austenite grain boundaries with a thin film of the molten metal. Usually, very low tensile stresses are required to fail parts that are liquid metal embrittled. In general, three conditions are necessary for liquid metal embrittlement. First, the embrittling metal must be present, either externally as a coating or internally. Internal sources can include lead used to enhance machinability. Second, temperatures that the part is exposed to
Energy absorbed, ft-lbf (J)
120 (160) 100 (130) 80 (100) 60 (80) 40 (50) 20 (30)
must be high enough that the embrittling metal can melt. Lastly, tensile stresses must be present as externally applied or internal residual stresses. Should any of these conditions not be met, then it is unlikely that the steel will fail by liquid metal embrittlement. Liquid metal embrittlement has been known to embrittle gun tubes. In 1977, during the manufacture of a 105 mm M68 gun tube, lead was electroplated to the tube and used as a lubricant during the autofrettage process. During the postautofrettage thermal treatment, the lead melted and embrittled the gun tube. A complete transverse brittle failure occurred. The axial tensile residual stresses from the autofrettage process were adequate to completely fracture the tube, even though the hoop stresses were much greater (Ref 46). In another example, an ISO 8.8 low-alloy steel bolt that was electroplated with cadmium was used for an extended time at an elevated temperature of 230 C (455 F). The resulting failure showed intergranular fracture, with cadmium penetration along grain boundaries. This cadmium penetration was detected by x-ray diffraction (Fig. 36) (Ref 9). On April 28, 1997, United Flight 1210, a Boeing 737–222 equipped with Pratt and Whitney JT8D-7B engines, experienced an uncontained failure of the No. 2 engine (right side, facing forward) high-pressure compressor stage disk during takeoff. Takeoff was aborted, and the aircraft was evacuated. Only two passengers were slightly injured during evacuation. Postincident examination of the engine revealed that two-fragments (approximately 50 by 100 mm, or 2 by 4 in.) separated from the disk. Examination of the disk revealed a 100 mm (4 in.) circumferential fracture around the diameter, with three additional fractures emanating diagonally outward toward the rim. Also, cracks Table 1 Melting temperatures of metals known to embrittle high-strength steels Melting temperature
0
Fig. 35
Shift in ductile-brittle transition temperature curve to a higher temperature for AISI 3140 steel by holding at 500 and continuous cooling through the temper embrittlement critical range. Source: Ref 35 C
Metal
°C
°F
Mercury Gallium Indium Lithium Cadmium Tin Lead Zinc Antimony
39 29 156 180 321 232 327 419 642
38 85 313 356 610 449 620 787 1187
Source: Ref 45
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emanated radially outward from two tie-rod holes, one of which bisected the fracture at the diameter. No cracks were detected on any of the other high-pressure compressor disks in the engine. On further examination by the National Transportation Safety Board, the fractures were found to have large intergranular areas in the steel compressor disk (Ref 47). Solid molten cadmium was detected along the prior-austenite grain boundaries, indicative of liquid metal (cadmium) embrittlement. Inadequate nickel plating prestrike thickness was observed at the surface of the disk. The failed disk was plated by a trainee and inadequately plated with nickel. It was found that the nickel plating was approximately 0.003 mm (0.00012 in.) in thickness, which is below the Pratt and Whitney specified thickness of 0.015 to 0.02 mm (0.0006 to 0.0008 in.). This thickness was inadequate to prevent migration of cadmium into the steel grain boundaries. This was not the first time that liquid metal embrittlement occurred in a compressor disk. On July 23, 1990, the crew of a JT8D-9-equipped Boeing 737–100 reported that they heard a muffled explosion during climb, followed by a loss of rpm of the No. 1 (left) engine. The crew returned to Houston, Texas, without incident. Engine examination revealed a failure of a disk spacer due to liquid metal embrittlement (Ref 48). An AISI 4330V hook-point failed during field trials. This hook-point is used to grab the cable on aircraft carriers and arrest the forward movement of the aircraft during landing. Previous hook-points failed because of excessive hardness and high triaxial stresses during impact loading. These hook-points were evaluated, and
Fig. 36
discrepant parts were segregated. A series of parts were then retempered to the specified hardness of HRC 46 from HRC 51. The lug radius was enlarged, and a new HVOF coating was applied. During field trials, multiple hookpoints were identified by magnetic particle inspection as having cracks in the lug radius (Fig. 37). No through cracks were found. The cracks were exposed, and a narrow uniform region of intergranular fracture (approximately 200 mm) was observed (Fig. 38). Metallography indicated that the microstructure was quenched and tempered martensite, and the hardness was within the specification of 46 to 48 HRC. Metallography revealed that no decarburization or precipitates were found at the prior-austenite grain boundaries. Hydrogen analysis showed a concentration of 0.2 ppm of hydrogen. The low concentration of hydrogen and the rapid rate of loading eliminated hydrogen embrittlement as a cause of cracking. Auger analysis of the grain boundaries within the intergranular region showed the presence of cadmium at the grain boundaries. The concentration of cadmium also decreased as the grain boundaries were ion milled away. This analysis indicated that the fracture occurred because of liquid metal embrittlement or solid metal embrittlement. Solid metal embrittlement is similar to liquid metal embrittlement, except that temperatures are not high enough to cause melting of the cadmium. For cadmium, solid metal embrittlement can occur at temperatures above 230 C (450 F). A review of the planning showed that the work order release did not include removal of the cadmium plating prior to retempering of the hook-points. Tempering to bring the hookpoints to the proper hardness was above 320 C (610 F). Based on this, it was determined that
Liquid metal embrittlement of a low-alloy bolt plated with cadmium that failed during service. Cadmium was found to have penetrated at the grain boundaries due to service above 230 C. (a) Overall fracture surface. (b) SEM examination of fracture showing intergranular fracture. (c) X-ray diffraction spectrum at grain boundaries showing cadmium penetration
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Fig. 37
AISI 4330V cadmium-plated hook-point used to arrest landings of naval aircraft. Overall view of the part, showing location of cracks observed using nondestructive testing
the cracking observed in the hook-points was due to liquid metal embrittlement caused by failure to remove the cadmium plating prior to tempering. Hydrogen embrittlement is a particularly insidious form of failure. Often, failure is delayed for hours, months, and possibly years after the component has been fabricated. The results may be catastrophic and unpredictable. The failure mode is typically intergranular along prior-austenite grain boundaries (Fig. 21). Hydrogen can come from either external or internal sources. One common source of
hydrogen is the steelmaking process (Ref 49), and it is a significant problem in large sections (Ref 50), where hydrogen embrittlement is observed as flakes or a reduction in ductility (Ref 51). These flakes or blisters are regions where hydrogen collects, until a bubble of hydrogen is adequate to deform the surrounding region. External sources of hydrogen are from manufacturing processes such as pickling (Ref 52, 53) and plating (Ref 54–56). Additional sources of hydrogen can be the result of galvanic coupling in an aqueous medium, in a similar fashion to electroplating. One of the particularly serious characteristics of hydrogen embrittlement is the incubation time required for it to occur. As a general rule, the higher the hydrogen concentration, the shorter the time to failure. For a given hydrogen concentration, as the stress is increased, the incubation time is decreased. In quenched and tempered steels, there are a number of sites that can trap hydrogen. These include martensite interlath interfaces, high density of dislocations, and the carbide-matrix interface. All of these sites can act as traps for hydrogen (Ref 57). Once present, hydrogen diffuses to traps, such as dislocation cores, and is transported by dislocation motion (Ref 58). Hydrogen can also collect at inclusions and carbides, which are also good hydrogen traps. The incubation time is dependent on the hydrogen diffusion rate in steel to the point of crack initiation. Quenched and tempered steels that have a hardness above HRC 38 are generally given a hydrogen embrittlement relief at 135 C (275 F) for 24 h. This enables the hydrogen in the part to diffuse out. This is based on the study by Johnson, Morlet, and Troiano (Fig. 39) (Ref 59). This hydrogen embrittlement relief is usually mandated whenever parts are plated, cleaned, or exposed in some fashion to aqueous solutions such as coolants or acid. Alkaline solutions are not generally prone to causing hydrogen embrittlement. Fasteners are prone to hydrogen embrittlement. In this example, an ISO 10.9 low-alloy steel bolt grade that was zinc plated failed during service. Multiple fracture initiation sites were evident along the bolt head transition, with intergranular fracture morphology and heavy secondary cracking. A hydrogen source was suggested from manufacturing (pickling stage) and/or cathodic hydrogen charging due to anodic zinc plating (Fig. 40) (Ref 9).
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To verify that the baking process after chromium plating was adequate, a plating shop tested four chromium-plated 4340 notched tensile
Fig. 38
specimens. These test specimens were heat treated to 1515 MPa (220 ksi) and chromium plated. During a sustained load test, one of the
SEM examination of the hook-point showing a narrow region of intergranular fracture along prior-austenite grain boundaries
Applied Stress, 1000 psi (MPa)
Normal Notch Strength = 300,000 psi (2070 MPa) (2070)
300
(1895)
275
(1725)
250
(1550)
225
(1380)
200
(1205)
175
(1035)
150
(860)
125
(690)
100
(520)
75
(345)
50 0.01
Uncharged
Bake 24 hr Bake 18 hr
Bake 12 hr
Bake 17 hr Bake 3 hr Bake 0.5 hr
0.1
1
10
100
1000
Fracture Time, hrs
Fig. 39
Baking AISI 4340 steel at 300 F for different times, showing the effect of baking on the incubation of failure
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Fig. 40
Hydrogen embrittlement failure of an ISO 10.9 low-alloy steel bolt grade. (a) As-received bolt. (b) Multiple initiation sites with secondary cracks evident. (c) Intergranular fracture along prior-austenite grain boundaries
Fig. 41
As-received notched tensile specimen showing location of fracture. Tensile specimen was fabricated from 4340 steel, heat treated to 1515 MPa (220 ksi), and chromium plated.
specimens failed prematurely (Fig. 41). The fracture was located at the notch. The fracture surface (Fig. 42) was examined, and the origin showed a shiny, faceted surface. At the origin, the fracture was intergranular, while away from the origin, near the center of the fracture surface, the fracture mechanism was microvoid coalescence. Hydrogen analysis on the notched tensile specimen yielded an average hydrogen concentration of 12 ppm hydrogen. This is considered very high and is sufficient to cause hydrogen embrittlement. Metallography of the test specimen showed a normal quenched and tempered microstructure, typical of a steel heat treated to this hardness. During a routine wheel and tire change, a new jack pad for a military aircraft failed, causing an aircraft to drop prematurely onto the new wheel. No damage occurred to the aircraft. The jack pad was machined from 300M steel that was heat treated to HRC 54 to 55. The jack pad was chromium plated. The as-received jack pad (Fig. 43) was examined, and two fracture surfaces were identified (Fig. 44). These were identified as origins 1 and 2. Ridges emanated away from a distinct origin location on each of the fracture surfaces.
Evidence of light corrosion products was found at the fracture origin of origin 2. SEM examination of each of the origins (Fig. 45, 46) revealed that the fracture was intergranular. At a distance away from the origin, the fracture consisted of microvoid coalescence, consistent with rapid ductile rupture. A metallographic specimen was removed from the largest origin location and examined (Fig. 47). The microstructure of the steel was quenched and tempered martensite, typical of 300M heat treated to HRC 54 to 55. Chromium plating was found to be intact at the fracture origin. Hydrogen analysis conducted on the jack pad showed hydrogen concentrations of 4 and 6 ppm, which is considered adequate hydrogen to cause hydrogen embrittlement in 300M steel heat treated to this hardness. Based on this investigation, it was concluded that the jack pad most likely failed from hydrogen embrittlement. Stress-corrosion cracking is the attack of a material by the combined action of tensile stress on a part, either externally from an applied force or internally from residual stresses, and a specific corrosive environment. Common features are brittle fracture with little ductility, localized corrosive attack, and a specific environmentalalloy system. Failure by stress-corrosion cracking (SCC) is characterized by exposure to a specific chemical environment and the simultaneous application of a tensile stress. Without one or the other, SCC will not occur. Fine cracks can penetrate deeply into the part without obvious signs of attack. Impending failure can occur without warning. The applied tensile stresses can be from the service environment or from any of the
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(a)
Fig. 43
As-received jack pad showing the locations of the two distinct origins on the inside bore of the hole for a pressed-in pin
(b)
(c)
Fig. 42
Overall view of the fracture surface, showing location and results of SEM examinations. (a) Overall fracture surface and location of origin. (b) Intergranular fracture at the origin of cracking (location A). Original magnification: 1000 · . (b) Microvoid coalescence at location B
Fig. 44
Fracture surfaces of the jack pad showing location of the origins. Original magnification: 2 ·
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numerous sources of residual stresses from manufacturing (thermal processing, machining, grinding, surface finishing, fabrication, or assembly). The tensile stress is important in the rupture of any protective film during initiation and subsequent propagation of the crack. There appears to be a threshold tensile stress intensity, KISCC, below which SCC does not occur. This stress intensity is dependent on
the alloy, the heat treated condition, and the environment. The site of initiation of SCC may be microscopic. This could be from local differences in metal composition or stress concentrations. A pre-existing mechanical flaw or discontinuity may act as a stress raiser and serve as a site for SCC initiation. Stress corrosion cracking usually exhibits extensive branching and propagates in a direction
(a) (a)
(b)
(b)
(c)
(c)
Fig. 45
Fig. 46
SEM examination of origin 1. (a) Location of the fracture origin. Original magnification: 20 · . (b) Location A showing a region of intergranular fracture along prioraustenite grain boundaries. Original magnification: 1000 · . (c) Location B, at a distance away from origin 1, showing microvoid coalescence. Original magnification: 2000 ·
SEM examination of origin 2. (a) Location of the fracture origin. Original magnification: 100 · . (b) Location C showing a region of intergranular fracture along prior-austenite grain boundaries. Original magnification: 1000 · . (c) Location D, at a distance away from origin 2, showing microvoid coalescence. Original magnification: 2000 ·
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perpendicular to the tensile stresses contributing to propagation and initiation. However, this is not always the case. Structural steels exposed to agricultural ammonia may exhibit nonbranched cracking. Stress-corrosion cracking has several special characteristics that differentiate it from other forms of cracking:
Only certain specific environments for a specific alloy system cause SCC. There is no general pattern regarding the corroding environments or alloy systems. Pure metals are much less susceptible to SCC. Cathodic protection has been successful in preventing the initiation of SCC or in stopping the propagation of SCC. Addition of certain soluble salts effectively can “poison” the environment and either reduce or stop the propagation of SCC cracks. Certain metallurgical features, such as grain size, can influence the susceptibility of an alloy system to SCC attack.
Macroscopically, fractures produced by SCC show little ductility and nearly always appear brittle. The fracture surfaces usually contain regions that are identifiable as the crack initiation site, slow crack propagation, and final failure. The regions containing the slow propagation often contain corrosion products or are discolored. This region extends to the region of final fast fracture. However, this can also be misleading, because the fracture could have corroded before inspection, or the environment may not be conducive to straining the fracture.
Fig. 47
Micrograph showing quenched and tempered martensite, typical of 300M heat treated to HRC 54 to 55. Note that the chromium plating is intact.
It is often difficult to differentiate between SCC and hydrogen-induced damage solely from the fracture surface. Fractures of both types exhibit intergranular features and tend to follow prior-austenite grain boundaries. Metallography is important to determine if branched cracking has occurred. Even so, the absence of branched cracking may not preclude SCC. In general, the environment that the part was exposed to can be the deciding factor of whether it is SCC or hydrogen embrittlement (Fig. 48). Low-carbon steels generally become more susceptible to SCC as the carbon concentration increases. Decarburized steels and pure iron are resistant to SCC. Microstructure plays a greater role in susceptibility to SCC than does the alloying elements. High-alloy steels in a variety of environments show that the heat treated strength of the alloy is more important than strictly the concentration. Steels that have been heat treated to 1240 MPa (180 ksi) or higher are especially susceptible to SCC. Typical environments that can cause SCC in steels are shown in Table 2. Caustic cracking in boilers is a serious SCC problem and has caused many failures in steam boilers. These failures usually initiate in riveted and welded structures, where small leaks allow buildup of caustic soda and silica. Cracking is usually intergranular. Failures of this type have occurred with concentrations of NaOH as low as 5% in water. Failures take place when the operating temperature is in the range of 200 to 250 C (390 to 480 F). The concentration of NaOH needed to cause cracking initiation decreases as the temperature is increased. Cracking of low-carbon steels and low-alloy steels in nitrate solutions occurs in tubing and couplings in high-pressure condensate wells. Cracking in nitrate solutions is intergranular, following prior-austenite grain boundaries. Generally, acidic solutions cause this type of cracking. Raising the pH of the solution enhances resistance to SCC, while increasing the concentration of nitrate-containing solutions tends to increase the susceptibility to SCC. Carbon steel tanks containing ammonia have also developed leaks because of SCC. Both plain carbon steels and quenched and tempered steel plate have shown a susceptibility to SCC in ammonia. Failures occurred in ammonia mixed with air and carbon dioxide. The presence of water vapor delayed cracking. Halide-containing environments, such as seawater, are particularly severe for alloy steels
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Direction of advancing cracking into metal
σ
σ
σ
σ H+ + e
H
++
M
++
+
M M+
Anodic Stress Corrosion Cracking Time to Cracking
Region of Anodic Stress Corrosion Cracking
Region of Immunity
Anodic Current M
Fig. 48
M++ + 2e
Hydrogen Embrittlement
Region of Hydrogen Embrittlement
Cathodic Current 2e + 2H+ –2H
Schematic differentiation of anodic stress-corrosion cracking and cathodic hydrogen embrittlement
Table 2 Environments that produce stress-corrosion failures in carbon and low-alloy steels Medium
Type of fracture (a)
Aqueous chloride environments
I,T
Caustic solutions Nitrates
I I
HNO3 HCN Seacoast and industrial environments Water, humid air, and gas
I I I
I
Comments
Prevalent in high-strength steels heat treated to 1380 MPa (200 ksi) or greater Well known as caustic embrittlement Examples of bridge cable failures in ammonium nitrate or sodium nitrate solutions ... ... High-strength steels heat treated to 1380 MPa (200 ksi) or greater are especially prone. High-strength steels heat treated to 1380 MPa (200 ksi) or greater are especially prone.
(a) I, intergranular failure; T, transgranular failure
heat treated to above 1380 MPa (200 ksi). The use of cadmium plating, low-hydrogen practices, and adequate baking are helpful
in preventing SCC in steels such as 300M or 4340. On August 22, 2003, an empty cargo tanker pulled upto a tank containing anhydrous ammonia. Approximately 1 hour after being filled, the front head cracked open (Fig. 49) and started to release anhydrous ammonia. Approximately 100 workers were evacuated from the building. Five people were treated for inhalation injuries and released. The cost to repair the trailer was approximately $25,000. Examination revealed a 40 cm (16 in.) long through-wall crack next to the radial weld in the front head at the 1 o’clock position (Fig. 50). Internal examination using magnetic particle inspection found two additional cracks that had not yet propagated through the wall of the tanker. SEM examination of the cracks (Ref 60) found that the fracture was branched and intergranular, with extensive surface corrosion on the crack faces.
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In the 1950s, the Agricultural Ammonia Institute determined that caustic cracking of ammonia-containing tanks was the reason that a number of carbon steel tanks had failed (Ref 61). They further determined that the addition of 0.1% water to anhydrous ammonia inhibited SCC in carbon steel. The committee recommended that at least 0.2% water be added to inhibit cracking. Further cracking occurred in the 1960s in quenched and tempered ASTM A517 steel, because purity levels had increased and water was no longer being added. In 1975, the Department of Transportation adopted regulations (Title 49 Code of Federal Regulations, Parts 171 to 180) that required cargo tanks fabricated from quenched and tempered steel should only be used for anhydrous ammonia if the solution contained 0.2% water. The
Fig. 49
regulation further required tankers to be placarded with signs indicating “QT” or “NQT,” for quenched and tempered or not quenched and tempered. The National Transportation Safety Board determined that the failure of the tank and the subsequent release of anhydrous ammonia were due to caustic cracking (SCC) of the tank from the transport of anhydrous ammonia containing less than 0.2% water. A Boeing 757–2008 was parked at a gate at Copenhagen, Denmark, and boarding of passengers was nearly completed when the righthand main landing gear truck beam failed (Ref 62). As the beam failed, the right side of the aircraft rested on the shock strut instead of on the wheels. Figure 51 shows the failed truck beam and the aircraft resting on the shock strut. A
Accident cargo tank with “QT” designation, which indicates quenched and tempered steel
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sketch of the main landing gear assembly of a B-757 is shown in Fig. 52. The fracture surface, showing evidence of corrosion, is shown in Fig. 53. Metallurgical analysis indicated that the fracture mode was due to SCC. Examination of the finish on the inner diameter showed that the plating on the inside diameter was thin or nonexistent, and that it did not receive the required shot peening. Because the truck beam was overhauled and the original plating was retained and worn during service, it was likely that the overhaul was inadequate or improper. This had the result of minimal cadmium protection on the inner diameter surface of the truck beam. Subsequent loss of the plating led to premature and severe corrosion in service and eventual fracture due to SCC. Creep Rupture The effects of temperature on mechanical properties and material behavior
are commonplace in everyday living. Examples include pipes bursting in the middle of winter, the expansion of a bridge in the middle of summer, and the sagging of a fireplace grate. Each of these examples is an indication that properties change with temperature. In addition, the previous discussion indicated that steels become more brittle as the temperature is decreased. There are many other effects of temperature that have been cited (Ref 63). Even the concept of elevated temperature is relative (Ref 64). What is considered hot for one material may be considered cold for another; for example, gallium has a melting point of 30 C, while tungsten has a melting point of approximately 3400 C. Creep is the continuous deformation of a material as a function of time and temperature. This topic is treated very thoroughly in Ref 65. The creep of a material is shown in Fig. 54. It can be seen from the figure that creep in a material occurs in three stages:
Stage I, where a rapid creep rate is seen at the onset of load application, then gradually decreases Stage II, where creep remains at a steadystate rate Stage III, where the creep rate shows an increasing rate until failure occurs
The behavior and creep rate are sensitive to the temperature to which the material is
Fig. 50
Through-the-wall crack on accident tanker
Fig. 51
Boeing 757-2008 truck beam failure occurring on Icelandic Air, aircraft registration TF-FIJ.
Fig. 52
Schematic of the assembly of a Boeing 757 main landing gear showing the location of fracture
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exposed, the surrounding atmosphere, and the prior strain history. Andrade and Chalmers (Ref 66) were pioneers in the study of creep and proposed that creep followed the equation: e=e0 (1+bt1=3 )ekt
where b and k are material constants that can be evaluated by several different methods (Ref 67). A better fit for the creep of materials was proposed by Garofalo (Ref 68). He indicated that:
testing, the time to failure is measured at a constant stress and constant temperature. This test has gained acceptance for elevated-temperature testing of turbine blade materials in jet engines. Using a tensile machine and high-temperature furnace, the strain is measured in creep testing by special extensometers suited for elevated temperatures. In stress-rupture testing, a simple apparatus such as a dial calipers is used, since only the overall strain at constant time and temperature is needed. Fatigue
de e=e0 +et (17e7n )+ t dt
where de/dt is the steady-state creep rate, e0 is the strain on loading, n is the ratio of the transient creep rate to the transient creep strain, and et is the transient creep strain. Very early, it was recognized that fractures at elevated temperatures occurred along grain facets (Ref 69). In stage III creep, intergranular wedge cracks and cavities form. Wedge-shaped cracks and creep cavities usually initiate at or near grain-boundary triple points and propagate along grain boundaries normal to the applied tensile stress. Creep cavities form at higher temperatures and lower working stresses. These structural features are shown in Fig. 55. Creep testing is usually performed for 1000 to 10,000 h with strains of up to 0.5%. Stressrupture testing, or testing to failure, uses much higher loads and temperatures, and the test is usually terminated after 1000 h. In stress-rupture
Fig. 53
Fracture surface of Boeing 757 main landing gear truck beam on Icelandic Air aircraft TF-FIJ
Parts are subject to varying stresses during service. These stresses are often in the form of repeated or cyclic loading. After enough applications of load or stress, the components fail at stresses significantly less than their yield strength. Fatigue is a measure of the decrease in resistance to repeated stresses. Fatigue failures appear brittle, with no gross deformation. The fracture surface is usually normal to the main principal tensile stress. Fatigue failures are recognized by the appearance of a smooth, rubbed type of surface, generally in a semicircular pattern. The progress of the fracture (and crack propagation) is generally suggested by beach marks. This is illustrated in Fig. 56 and 57. The initiation site of fatigue failures is generally at a stress-concentration site or stress raiser. A typical fracture appearance is shown schematically in Fig. 58. Three factors are necessary for fatigue to occur. First, the stress must be high enough that a crack is initiated. Second, the variation in the stress application must be large enough that the crack can propagate. Third, the number of stress applications must be sufficiently large that the crack can propagate a significant distance. The
Fig. 54
Schematic representation of creep
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fatigue life of a component is affected by a number of variables, including stress concentration, corrosion, temperature, microstructure, residual stresses, and combined stresses. The structural features of fatigue failures are generally divided into four distinct areas (Ref 70):
Crack initiation, the early development of fatigue damage Slip band crack growth, the early stages of crack propagation. This is often called stage I crack growth. Stable crack growth, which is usually normal to the applied tensile stress. This is called stage II crack growth. Unstable crack growth, with final failure from overload. This is called stage III crack growth.
Fig. 55
Fig. 56
Creep cavities and creep wedges forming at grain boundaries
Actual fatigue failure of a crankshaft showing characteristic beach marks. Fatigue initiated at the radius of the journal and exhibits classic bending fatigue.
Fatigue usually occurs at a free surface, with the initial features of stage I growth, fatigue cracks, being initiated at slip band extrusions and intrusions (Ref 71, 72). Cottrell and Hull (Ref 73) proposed a mechanism for the formation of these extrusions and intrusions (shown schematically in Fig. 59) that depends on the presence of slip, with slip systems at 45 angles to each other operating sequentially on loading and unloading. Wood (Ref 74) suggested that the formation of the intrusions and extrusions was the result of fine slip and buildup of notches (Fig. 60). The notch created on a microscopic scale would be the initiation site of stable fatigue crack growth. In stage II, stable fatigue crack growth, striations (Fig. 61) often show the successive position of the crack front at each cycle of stress. Fatigue striations are usually detected using electron microscopy and are visual evidence that fatigue occurred. However, the absence of fatigue striations does not preclude the occurrence of fatigue. Striations are formed by a plastic blunting process (Ref 75). At the end of the stage I crack tip, there exists sharp notches due to the presence of slip. These sharp notches cause stress to be concentrated at the crack tip. The application of a tensile load opens the crack along slip planes by plastic shearing, eventually blunting the crack tip. When the load is released, the slip direction reverses, and the crack tip is compressed and sharpened. This provides a sharp notch at the new crack tip where propagation
Fig. 57
Fatigue failure of a fastener, with initiation of fatigue occurring at the threads
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Fig. 58
Schematic illustration of simple fatigue failures
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can occur. This is shown schematically in Fig. 62. An alternative hypothesis on striation formation was presented by Forsyth and Ryder (Ref 76). In their model, the triaxial stress state at the crack tip forms a dimple ahead of the crack front. The material between the crack tip and the dimple contracts and eventually ruptures, forming a fatigue striation. This is shown schematically in Fig. 63. In mild steel, well-defined striations are observed but not as well defined or as spectacular as in aluminum. This was first assumed to be due to the crystal lattice structure, since face-centered cubic austenitic steels show welldefined striations, and mild steels (basestructured) do not (Ref 77). Other alloys, such as titanium alloys, with a hexagonal close-packed crystal structure show very defined striations (Ref 78). However, aluminum alloys (bodycentered cubic) show strongly defined striations (Ref 79). Therefore, attributing defined striations to crystal lattice alone was discounted as a viable theory. Deformation and available slip systems were presumed to be more significant (Ref 80). However, this does not follow, because mild carbon steels are more ductile than austenitic steels. It is now generally accepted that
fatigue striations form by the plastic blunting process. It has also been found that the thicker the testpiece, the faster the crack propagation rate (Ref 81). It is likely that the propagation rates for thicker pieces are due to increased plane-strain conditions, with a small plastic zone at the crack tip. Since there is a greater stress gradient for a small plastic zone, a faster crack propagation rate may be expected. Also, in thicker panels there is a higher state of triaxial stress, which would also tend to increase crack growth rates. Since fatigue failures usually begin at the surface, the surface condition is very important. Surface roughness is a primary factor influencing fatigue. Highly polished specimens exhibit the longest fatigue life, with increasingly rougher surfaces yielding decreased fatigue life. Rough lathe or coarse grinding reduces the fatigue strength by approximately 20% below polished specimens (Ref 82). Electropolished specimens have lower fatigue limits than mechanically polished specimens, by up to 25% (Ref 83). This reduction is due to the removal of surface compressive residual layers induced during mechanical finishing. An example of a typical fatigue failure in an ASTM B7 low-alloy steel bolt grade is shown in Fig. 64 (Ref 9). Fracture initiation occurred along the threads with typical and pronounced beach marks (i.e., cyclic fracture propagation) and transgranular fracture mode. An example of a manufacturing effect on fatigue is the following example of an arresting gear hook shank (Fig. 65) used to slow down aircraft when landing on aircraft carriers. In this example, the hook failed after 1361 simulated arrestments, which was below the lifetime of 2250 arrestments. The part is designed to last
Fig. 59
Schematic representation of the mechanism of fatigue intrusions and extrusions
Fig. 60
Mechanism of intrusions and extrusions. PSB, persistent slip bands. Source: Ref 74
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two lifetimes, or 4500 arrestments, without cracking (0.25 mm, or 0.010 in. detectible flaw). The arresting shank was fatigue tested in a fixture, with hydraulic cylinders providing loads at the vertical damper and hook-point cable groove. The maximum applied load was 90 mg (200,000 lb). A schematic of the arresting hook shank is shown in Fig. 65. The arresting hook shank was fabricated from an AerMet 100 rotary forging. It is rough turned on a lathe on the outside, then gun drilled to create a pilot hole down the length of the forging. The outer surface is turned to the final
Fig. 61
Typical fatigue striations in 7075 aluminum
Fig. 62
Mechanism for fatigue striation formation
diameter. The bore is then injection drilled to the final dimensions. A follower supported the injection drill. This is not a method that is commonly used for final machining operations. It is heat treated in vacuum to 1930 MPa (280 ksi) ultimate tensile strength. The part is inspected using dye-penetrant and magnetic particle nondestructive testing methods. The bore is visually inspected using a bore scope. This is a difficult inspection because of the long length and narrow bore. Examination of the fracture surface showed that cracking initiated at the hook-point side, on the inner diameter, at a location approximately 26.5 cm (10.5 in.) aft of the uplock retainer. The fracture had characteristics of fatigue fracture, with multiple origins observed. Surface roughness measurements varied across the inner bore, from approximately 1 to 5 mm (40 to 180 min.). The drawing requirement was 3 mm (125 min.). Circumferential machining marks were found at the fracture origin (Fig. 66). SEM examination (Fig. 67) showed fatigue striations emanating away from the identified origin. Cracking was found to have initiated at circumferential machining marks. Machining marks were observed at 4.3 mm (0.17 in.) intervals. Many secondary cracks were observed at the machining marks. Fatigue was found to initiate subsurface to the inner bore, adjacent to the machining marks. A welldefined surface layer was observed. This layer had the appearance of mechanical working or damage. This observed layer followed the feeds and speeds of the injection drill. Metallography showed that the material was quenched and tempered martensite and was typical for this material heat treated to this hardness. At the
Fig. 63
Striation formation from ductile dimple formation ahead of a crack front
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Fig. 64
ASTM B7 low-alloy steel bolt grade. Fracture initiated along threads, with typical and pronounced beach marks (i.e., cyclic fracture propagation) and transgranular fracture mode. (a) Location of bolts in pump coupling. (b) Beach marks showing asymmetrical bending with initiation at high stress-concentration factor at bolt threads. (c) Transgranular fracture morphology
Fig. 65
Fig. 66
Schematic of the failed arresting hook shank showing location of loads
Machining marks found on the inside of the bore, at the origin of cracking
origin, the presence of well-defined subsurface cracking was observed. This layer had the appearance of smeared metal and base metal pullout. Flat cracking, suggestive of fatigue cracking, was observed to emanate from the flaw (Fig. 68). The flaws were located at 4.3 mm (0.017 in.) intervals, identical to the feed rate of the injection drilling process. During the injection drilling process, three cutters are used. Coolant is forced through a central hole to cool the cutting tools and to flush the chips. AerMet 100 tends to form long strands of material during machining and does not want to form chips. Hot chips can contact the freshly machined surface. These chips or long strands are under pressure at the cutter or follower and can be forced onto the newly machined surface by the follower. If the temperature and pressures are high enough, solid-state welding of the chips and bore surface can occur. As the cutter boring bar moves, pullout can occur. The
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Fig. 68
Metallography of the arresting hook shank. (a) Typical quenched and tempered martensite found. This is typical for the hardness of the arresting hook shank. (b) Pulled material at 4.3 mm (0.17 in.) intervals along the inner bore of the arresting hook shank. Origin is to the left. (c) Secondary cracking observed at the location of pulled material
Fig. 67
SEM examination of the fracture surface. (a) Fatigue striations emanating from the fracture origin. (b) Machining marks found on the surface of the inner bore. (c) Welldefined layer showing fatigue emanating from the damaged material at the surface of the inner bore
examined flaws matched the machining feeds and speeds. The arresting hook failed by fatigue, initiating at flaws created during the final machining process. The defect morphology suggested localized solid-state welding and pullout from chip contact with the freshly machined surface. The surface roughness and finish of the inner bore did not meet drawing requirements.
Summary In this short overview of the possible mechanisms of failure for steels, the following were discussed:
Techniques for examining fractures Ductile and brittle failures Intergranular failure mechanisms Fatigue
The previous discussion has shown that it is important to look at not only the fracture surface but at all the factors (manufacturing history, service conditions, and loading). All the tools available to the metallurgist should be used— these include photography, fractography, and
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metallography—to understand the sources and root cause of failure.
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31. J.F. Knott, Fundamentals of Fracture Mechanics, Wiley, New York, 1973 32. A.H. Cottrell, Proc. R. Soc. A, Vol 276 (No. 1), 1963 33. A.H. Knott, Mater. Sci. Eng., Vol 7 (No. 1), 1971 34. J.S. Colton, Class notes, “Great Boston Molasses Disaster, Jan 15, 1919,” ME6222 Manufacturing Processes and Systems, Georgia Institute of Technology 35. G. Krauss, Steels: Heat Treatment and Processing Principles, ASM International, 1990, p 231, 237 36. G. Thomas, Retained Austenite and Tempered Martensite Embrittlement, Metall. Trans. A, Vol 10, 1979, p 1643–1651 37. S. Banerji, C. McMahon, and H. Feng, Intergranular Fracture in 4340Type Steels: Effects of Impurities and Hydrogen, Metall. Trans. A, Vol 9, 1978, p 237–247 38. D. Kalderon, Steam Turbine Failure at Hinkley Point ‘A’, Proc. Inst. Mech. Eng., Vol 186, 1972, p 341–377 39. L.D. Kramer and D. Randolph, Analysis of TVA Gallatin No. 2 Rotor Burst: Part 1—Metallurgical Considerations, Proc. 1976 ASME-MPC Symposium on CreepFatigue Interaction, p 1–24 40. H.K.D.H. Bhadeshia, High Performance Bainitic Steels, Mater. Sci. Forum, Vol 500, 2005, p 63–74 41. I. Olefjord, Temper Embrittlement, Review 231, Int. Met. Rev., Vol 23, 1978, p 149–163 42. B. Woodfine, Temper Brittleness, A Critical Review of the Literature, JISI, Vol 173, 1953, p 229–240 43. F. Carr, M. Golman, L. Jaffee, and D. Buffum, Isothermal Temper Embrittlement of SAE 3140 Steel, Trans. TMS-AIME, Vol 197, 1953, p 998 44. Liquid Metal Embrittlement, Failure Analysis and Prevention, Vol 10, Metals Handbook, 8th Ed., American Society for Metals, 1975, p 228–229 45. P. Fernandez, R. Clegg, and D. Jones, Failure by Liquid Metal Induced Embrittlement, Eng. Fail. Anal., Vol 1 (No. 1), 1994, p 51–63 46. G. Vilganlte, G. Troiano, and C. Mossey, “Liquid Metal Embrittlement of ASTM A723 Gun Steel by Indium and Gallium,” ARCCB-TR-99011, Army Research, Development and Development Center,
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73. A.H. Cottrel and D. Hull, Proc. R. Soc. (London) A, Vol 242A, 1953, p 211 74. W.A. Wood, Bull. Inst. Met., Vol 3, 1955, p5 75. C. Laird, “Fatigue Crack Propagation,” in STP 415, American Society for Testing and Materials, 1967, p 136 76. P.J. Forsyth and D.A. Ryder, Metallurgica, Vol 63, 1961, p 117 77. G. Jacoby, Current Aeronautical Fatigue Problems, J. Schijve, Ed., Pergamon, New York, 1965, p 78 78. W.R.Warke and J.M. McCall, “Fractography Using the Electron Microscope,” ASM Technical Report We-2-65, American Society for Metals, 1965 79. G. Jacoby, Fractographic Methods, Exp. Mech., 1965, p 65 80. P.J. Forsyth, A Two Stage Process of Fatigue Crack Growth, Symp. Crack Propagation, Vol 2, Cranfield, U.K., 1961, p 76 81. D. Broek and J. Schijve, “The Influence of Sheet Thickness in the Fatigue Crack Propagation in 2024-T3 Alclad Sheet Material,” NLR Technical Report M2129, Amsterdam, 1963 82. N.E. Frost, K.J. Marsh, and L.D. Pook, Metal Fatigue, Oxford University Press, London, 1974 83. T.T. Oberg and E.J. Wad, Technical Note Report 56-289, Wright Air Development Department, 1956
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Failure Analysis of Heat Treated Steel Components L.C.F. Canale, R.A. Mesquita, and G.E. Totten, editors, p 87-109 DOI: 10.1361/faht2008p087
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
Mechanisms and Causes of Failures in Heat Treated Steel Parts Debbie Aliya, Aliya Analytical, Inc.
THE TERM mechanism of failure means different things to different people. One possible definition refers to a particular product application, a particular type of component, and a particular type of industry where certain environmental conditions are common. In many cases, these mechanisms may apply only to a limited type of material or an alloy family. In order to understand particular failure mechanisms, it is important to understand the causes of failure and categories of damage. This chapter reviews various ways to classify failure categories. Information on mechanisms of damage, in particular environment-material-type pairs, is pro-vided toward the end of this chapter. The term damage is often preferable to the term failure, because damage is a technical term that is very clear and has a specific meaning, generally related to the physical condition of the component. The term failure, on the other hand, has many more philosophical connotations. It is possible to see at once, perhaps, that the part is damaged. Only at the end of the investigation does one have a good chance of knowing with a high degree of engineering certainty whether the part itself failed, rather than the design, the design system, the employee training system, the procurement system, or the user-certification system. Thus, it is important to avoid using the term failed part until one has truly determined that the part was the problem. Often, one of the important goals in a failure analysis of a heat treated part is to determine whether the damage is the result of improper heat treatment, that is, the heat treater’s fault. For many people in manufacturing, especially if an independent or “job shop” heat treater is involved, it is easy to blame the heat treater. After all, he was the last one to touch the part before assembly, and most people in the general manufacturing arena understand heat treating poorly, if at all.
If the part does not meet the specification for mechanical or physical characteristics after heat treating, it may be the heat treater’s fault. However, there may have been something wrong with the raw material, or prior manufacturing processes that allowed a part that went through the normal heat treating and inspection process to have substandard properties, which would not be the heat treater’s fault after all. This chapter gives some examples of lack of conformance to specification that may at first look like the heat treater did something wrong, but where other contributing factors made it difficult or impossible for the heat treater to meet the specification. This chapter also summarizes the basic types of damage, with particular consideration given to whether their likelihood can or cannot be influenced by the heat treating process. The classical organization for types of damage (failures) is as follows: deformation, fracture, wear, corrosion or other environmental damage, and multiple or complex damage. Separately from what the damage type is, one also should look at the potential causes, sources, or factors promoting the damage. Failure analysts used to be taught to classify failures as a result of defects or abuse. There is still a large amount of literature that presents failure causes in such binary terms. By limiting the analysis to one of two possible causes, opportunities may be missed for improving the product. There are actually several different ways to classify causes of failures and damage. This chapter also describes a process that can be used to demonstrate likelihood that a product was abused. Sometimes, the physical evidence speaks clearly. If this is not the case, it may be necessary to quantitatively prove, for example, that the part was overloaded. This may be difficult to do, because exact service conditions are
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frequently impossible to determine with any degree of certainty. Besides service or maintenance abuse, four potential sources or origins of damage and failure relate to the product life cycle, as follows:
Design process omissions Undesired raw material characteristics Undesired component manufacturing characteristics (includes heat treating) Improper service or maintenance conditions
In this chapter, a modified classification is developed specifically for failure analysis of heat treated parts where the heat treating is suspected to be the cause of the failure. The first logical potential source of poor component performance is raw material characteristics. Note the use of the word characteristics rather than defects. This is an important thing to keep in mind when performing this type of work. The term defect has a specific legal definition that does not necessarily apply to all cases of physical flaws or suboptimal material properties. Thus, the use of the word defect may have undesirable cascading consequences, especially when personal injury or large financial losses were a result of the component malfunction. The second potential source of poor performance is undesirable component characteristics. For example, castings, forgings, and machined or molded components may have discontinuities or microstructural features present that make it appear to the casual observer that there was a heat treating problem. These discontinuities can contribute to a poor heat treating outcome, a poor service outcome, or both. The third potential trouble spot is design characteristics, many of which may fall under the subclassification of inadequate attention to detail. The fourth category of causes of failure in this scheme is true heat treating process problems. Note that in this classification, these are, so far, all things that can go wrong in the engineering and manufacturing of the component. Finally, service and/or maintenance abuse can be considered.
cannot be created or facilitated as a result of a problem in the heat treating process. Deformation Figure 1 shows the result of a major deformation event. Note the image of a cylindrical structure. There is a dark sign-wave-shaped band where the material is crumpled. The weight of the structure above that area created the force that caused the cylinder to suddenly deform. There were also some tears or fractures as well, but the main visible damage type is deformation. There are different ways to categorize deformation. One way is to compare gradual to sudden deformation. Gradual deformation can occur when something is loaded and the load is sustained. Due to the sustained load, the structure can stretch or bend. This type of deformation can happen during heat treating. Imagine a part that has a protrusion that is not supported by an appropriate fixture. The part is heated to red heat. Depending on the particular configuration and the presence of residual stresses in the component, the protruding feature may droop due to gravity, or it may change its shape in some other way to relieve residual stresses from earlier parts of the processing. It is important to consider this type of deformation in a component that has a critical dimensional envelope. If the part goes outside of the specified or required envelope and is placed into service, the stresses that are experienced by the dimensionally nonconforming component may be quite different from the original design intent. Gradual onset deformation may also happen in parts that are loaded near the elastic limit. It is important to note that the published elastic limit for different materials is a numerical quantity that is determined by using a specific set of test parameters, usually of relatively short duration. Long times at a stress level that may be
Types of Damage and Failure The four basic types of damage (deformation, fracture, wear, and corrosion) are briefly reviewed in this section, and discussions address whether each of the four types of damage can or
Fig. 1
Example of a sudden deformation event due to buckling
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considered to be below the yield strength may actually cause unacceptable permanent deformation. Since the yield strength or elastic limit can be affected by heat treating, it is possible, although probably not a frequent cause, that this type of service-related gradual onset deformation is also related to a heat treating problem. Gradual onset deformation could also be a result of loading the part beyond its design intent or actual physical limitations. The alternative to gradual onset of deformation is sudden onset. Examples of sudden onset deformation are buckling instability in compression of columns or torsion of tubes. A roof truss that collapses under a snow load is such an example. Roof trusses are not usually heat treated for strength, at least not separately from the thermal element of the hot rolling. However, it is possible that annealed steel may be used in some structural components. Most sudden onset damage is primarily related to the basic geometry and modulus of elasticity, which is not a strong function of any heat treating process. Thus, heat treating problems are generally a minor or insignificant factor in most sudden onset damage events. Another way to classify deformation is by level or degree: elastic (in other words, if the load is removed, deformation is relieved) or plastic (which is permanent deformation). A possible example of an elastic deformation failure is a spring that does not have the correct spring constant. Imagine a coil spring that is supposed to stretch out 0.1 mm when subjected to 12 N of force. What if the spring stretches out 0.05 or 0.2 mm? Can that type of failure be due to heat treating? What are the causes or the factors that allow these two types and two levels of deformation to happen? As with sudden onset, the two main factors that control elastic deformation are geometry and modulus of elasticity. These are not factors that are greatly influenced by the heat treater. Thus, the heat treater is usually not at fault in the case of sudden onset buckling of a column, sudden onset buckling of a tube in torsion, or for elastic deformation failures. In general, sudden onset and elastic failures are a result of the combination of the design or actual geometry and the elastic modulus being insufficient to sustain the loading conditions. Sudden onset and elastic deformation damage is only secondarily related to the yield strength. The modulus of elasticity is, in general, a constant. However, modulus is not totally a constant in wrought
materials, where differences can occur due to preferred crystal orientation. This is less likely in heat treated steel than in steel used in the asworked condition. In general, in most heat treated steels, the modulus of elasticity will vary little. In attempting to determine the predominant factor or predominant cause in a sudden onset deformation, one should look first at the geometry and the applied loading. Heat treating issues would be considered in puzzling cases that did not lend themselves to ready analysis. For gradual onset deformation, the geometry and the elastic modulus are still important, but the heat treating can have a much more significant effect. If the yield strength is not as high as it is supposed to be, unacceptable levels of gradual onset deformation may occur. This concludes the introduction on how to think about deformation failures and whether the heat treater even needs to be involved. Figure 1 shows what was reported to be a sudden onset event. This is actually an example of a complex failure or damage mode, because this structure was standing for quite a few years, and then one day, the wind blew and the damage happened. Investigators found evidence of long-term corrosion on the inside, associated with significant wall thinning in some areas. As another example, if a hollow tube for a truck drive shaft is not heat treated properly, could it buckle more readily? If the damage is buckling deformation, the primary factor would be the wall thickness and the modulus of elasticity. However, there may also be some strength issues. When people talk about buckling due to instability (or Euler’s buckling, or what is referred to as sudden onset), the main factors are the geometry and the modulus of elasticity. However, with geometrical configurations that are not exactly like the extreme examples that Euler used to develop his theory, one can appreciate that the equations become very complicated and actually do have factors based on strength values. However, the predominant factors, in general, are the geometry and the modulus. In failures of heat treated structural steel, the modulus will be 30 million psi, with some variations. For stainless steel, the modulus may be a little different. If the component material is a heavily cold-worked steel with oriented, heavily textured microstructure, then the modulus of elasticity may be different in the different directions. However, most of this orientation may be eliminated by any
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subsequent heat treating that included an austenitization phase. Note again that this discussion is strictly speaking about deformation in the absence of a primary fracture event. Before leaving the subject of deformation, consider a small spring clip that was designed to be used at the top of the linear portion of the stress-strain curve. Why would anyone design a spring to be used at a stress almost at the nominal yield strength? No experienced metallurgical engineer would expect each one of a quarter of a million such spring clips to sustain multiple load cycles to the theoretical elastic limit and not have any permanent shape change. These clips could not even sustain the minimum 15 loadcycle requirement without excessive permanent deformation. This was a case of a nominally elastic spring application where deformation failure was caused by the heat treatment, specifically by normal variations of the heat treating process. It would be difficult to blame the heat treater in the absence of very strict qualitycontrol specifications. In this case, the major cause of the failure was the design engineer’s unrealistic expectations of uniformity of heat treated components or the design process in the company in which the designer was working.
Fracture How can fractures of heat treated steel parts be examined to determine the existence of any factors related to the steel itself or its heat treating? Basically, there are two types of questions to ask. One of the things that people like to ask is, “Why did this one break?” That is a useful question in the case of a part that has been used successfully for many years. Maybe there is one part that failed out of a half-million parts that are in service. Proper examination and evaluation of the physical evidence can reveal much to answer this question. Before evaluating the effects of the heat treatment itself, one must first examine the physical condition of the damaged property. It must be understood how the loads interacted with the component to create new surface area where none used to exist. The visual appearance in three dimensions can reveal a large amount of information on issues related to how the part was really loaded. The colorations and surface texture of the newly formed (undesired) crack surfaces can indicate how the crack happened and how long it took for the crack to grow.
The other question that is frequently asked is, “Why did this one break at this specific time?” This is a different question, and to answer it, fracture mechanics type of explanations and theories must be explored. In such a case, one assumes that every structure has some small discontinuity and a related crack growth rate, which is a function of the stress intensity (measured in terms of the mathematical product of the nominal component stress and the crack size) and the fracture toughness. These are functions of service condition and material parameters. Particularly for parts that appear to have broken in fatigue, (consider a two-year service life when no similar component had previously cracked in under five years), it may be informative to look at the microstructure and how that may have impacted the fracture toughness. Microstructure and fracture toughness could definitely be related to heat treating issues. Other chapters in this book give more information on fracture mechanics. For a more conceptual, lower-math-content methodology to understand why something cracked or why it broke, start by reviewing the stress and strength variations that are at work in the component. Any place where the local stress exceeds the local strength can initiate a crack. There can even be single grains that are low strength for some reason. If the local strength is lower than the local stress, then it is possible to initiate a crack. Once a crack exists, it may or may not progress to complete component fragmentation. Sometimes, cracks do not propagate. However, cracks generally do not heal themselves, and they often do propagate. It is actually reasonably straightforward to learn to look at a component, if one has an understanding of the loading geometry and the heat treating, and determine if there was something wrong with the component in question. Someone inspecting a component with a complex shape or complicated loading history may benefit from a good finite element analysis. However, many components do lend themselves to ready evaluation regarding the presence or absence of an extraordinary factor promoting premature fracture. To use the stress/strength distribution concept to analyze fractures, one should be familiar with the six basic loading geometries, including tension, compression, bending, torsion, contact stress, and shear. A review of some of the loading geometries is presented in Ref 1; however, the three-dimensional characteristics of the
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fragments are not systematically presented. Armed with information on these basic fundamental loading geometries, one can learn to predict macroscale fracture appearance in ductile and brittle materials given different loading conditions, along with the more well-known fracture surface features, such as beach marks, ratchet marks, and chevrons. This is a much more powerful method of visual examination than simply interpreting surface texture features. With practice, one can look at a broken component and obtain a realistic idea of what loading conditions actually caused the fracture. Sometimes, it can be shown that the real loading was very different from the design loading. Before leaving this subject, a review is needed on details that are often poorly understood relating to how to distinguish ductile and brittle cracks on the macroscale. This is clearly very important for a heat treater to know to defend against incorrect accusations of embrittlement. There are some types of macroscale ductile cracks that can easily be misinterpreted as brittle
cracks. In fact, much of the published literature is unclear on this issue. To understand this more clearly, see Fig. 2. The image is of a broken tensile bar. Near the fracture, the material is necked down. Because of the visible shape change, this is an obvious ductile fracture. The image in Fig. 3 is of a threaded fastener. Based on the ridge patterns in evidence, the crack started at the root of one of the threads and went back into the page. There is very little indication of any deformation visible at the macroscale. This is correctly called a macroscale brittle fracture. Figure 4 has more challenging fragments for fracture analysis. The image shows a chain link, which did not originally have an open shape. The two protruding ends were touching; there was no gap. The lower portion of the link did not have a curved arc shape as depicted. A significant amount of deformation was clearly associated with this fracture event. Does that make this a
Fig. 2
Fig. 3
Example of a macroscale ductile fracture in tensile loading
Example of a macroscale brittle fracture in tensile loading
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ductile fracture, or could this be a brittle fracture? Imagine that only this one photograph is available, and it must be determined whether this was a ductile or brittle crack. Clearly, the steel that was used to make the link has the capability of being deformed. In other words, the steel is ductile at the conditions present at the time of bending. The analyst must not be satisfied with that answer, though. The proper analysis includes a determination of when the deformation happened. Postfracture deformation does not make the crack event itself ductile. The macroscale deformation must occur during and as an inseparable part of the fracture process for the crack itself to be a ductile crack event. For those unfamiliar with this methodology, imagining that the link broke suddenly while it was under load can be helpful. If the crack were ductile, one would likely see some necking at the crack location, since chain links are generally loaded in tension. No such localized deformation is observed in this case. On closer examination, one can see a tiny shear lip at the top edge of the fracture surface. That makes this crack, for the most part, a brittle crack at the macroscale, despite the presence of available ductility in the material. The visible deformation near the center of the lower (originally straight) portion of the link happened after the crack was completely formed. In doing fracture analysis, it is important to distinguish the capability of ductility in the material from the behavior at the time of the crack event. Despite the material ductility, the crack happened in a brittle way. Closer examination of other views not shown provide clear evidence that this was a fatigue crack. Beach marks were visible. Fatigue cracks grew below the yield strength of the component, creating macroscale brittle features. To clarify one other potential source of confusion, it is important to remember that tensile refers to a loading geometry. Fatigue is a type of crack
Fig. 4
Example of a macroscale brittle fracture in tensile loading
path. In this case, a fatigue crack grew due to a tensile load in the horizontal portions of the link (as shown). To add a few more details to this case study, the chain was in service at a plant that processes meat, and strong acids were used to clean the conveyor systems. This crack actually initiated at a corrosion pit on the inside surface. The cleaner reached the inner surfaces, but the employees may not have rinsed the chain very well. This allowed a corrosion pit to form, which then allowed a fatigue crack to grow. Again, it is important to understand that the material itself is ductile; there is nothing wrong with the material. People involved with failure analysis need to keep in mind that material behavior is or at least may be different from material capability. People doing failure analysis work need to be able to distinguish inherent capability and actual behavior. To underscore the importance of separating the behavior from the capability, imagine the potential corrections that may be considered if someone found this to be a ductile overload fracture. The “cure” may be to make it harder. In the case of the acid cleaning, harder steels are often more susceptible to stress corrosion than softer steels. If the “harder-is-the-answer” theory were put into practice, an undesirably short life may become a horribly short life. It is important to be sure that a crack that is diagnosed as ductile is really ductile and one that is diagnosed as brittle is really brittle. Finally, returning to the heat treating issues, the fact that someone misdiagnosed this crack as a ductile fracture may lead to the heat treater being blamed for overtempering or inadequate hardening. In fact, until now, the potential blame or innocence of the heat treater has not been investigated in any thorough manner. It is possible that poor heat treating or poor material manufacture contributed to the ease of corrosion attack, and the meat processing plant employees were blamed incorrectly. Further examination of the microstructure is required to reveal the root physical cause of the fracture and its timing. Finally, there is the shaft fragment shown in Fig. 5. The shape is cylindrical. The image shows one fracture face. No necking or reduction in area is visible at the fracture face location. However, this is a ductile fracture. It is necessary to know that this shaft broke in torsional loading. In torsion, the shear stresses are in the transverse orientation to the length of the shaft. To best understand macroscale ductile and
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brittle fracture, one must be familiar with normal and shear stresses. On a simple level, normal stresses cause macrobrittle cracks due to crack opening forces created at the crack tip. Shear stresses allow slip and are the basis of the deformation that creates the ductile crack event. So, even though there is no necking, this is a ductile crack. Often, it is relatively easy to see some evidence of twisting on the side of a ground shaft that has failed from ductile fracture by torsional forces, which would lend more credibility to this diagnosis of ductile fracture. However, this shaft was extremely smooth, and it required a long etching time in a heated acid solution to reveal permanent twisting on the original cylindrical surface. Another possible source of confirmation that this is a ductile crack is the classical smear features on the fracture face. It has been argued that these smeared features could be a result of postfracture damage. While this is a possibility that should be considered, since the background evaluation revealed that the shaft was loaded in torsion, and there are no crack opening stresses operating on the transverse planes, this must be a macroductile crack. A macrobrittle crack in torsion is helical. In closing this section on fracture, note how important it is to follow the advice of the many authors and teachers who state that background research is step one of a failure analysis. The
Fig. 5
Example of a macroscale ductile fracture in torsional loading
possible loading geometries that could have created the fracture must be reviewed to make a proper determination of whether or not the crack is macroscale ductile or brittle. That determination cannot be made without assessing what the loading geometries may have been. If this fragment had been totally covered with red rust, it would have been even more difficult to determine the basic ductile or brittle behavior of the material without knowledge of the loading geometries and expected fragment shape. Stress versus Strength. Almost all real loading geometries cause the stress to be highest somewhere along the part surface. If the strength is uniform, for example, if there is a piece of hot rolled 1050 steel that does not have any decarburization or carburization and has not been shot peened, the crack initiation is expected somewhere at the original part surface. In the presence of any type of bending or torsional stresses, the highest stress will be at the surface of the part. In the presence of pure tensile loading, theoretically the crack could start anywhere in the cross section. Such pure tensile loading is rare. Imagine the case of a hydraulic cylinder rod. Even here, there must be a section change, a fillet, at some point. The loading at the fillet is not uniform; there is a stress concentration. Even a tensile test coupon that is forced to break in an area of nominally uniform strength and stress is not totally uniformly loaded. Most tensile test coupons are tapered so that the stress is slightly higher at the center of the gage length. This brings the discussion to what is so useful about heat treating steel. Many types of steels and heat treatments create harder or stronger layers at the surface. Heat treating allows the strength to be increased where it is useful. Figure 6 is from Ref 1with annotations. This figure shows a subsurface crack initiation along the boundary between the induction-hardened case and the softer core. A second initiation appears to be inside the induction-hardened case. It must be recalled that cracks can happen whenever the local stress is higher than the local strength. Normally, the stress is expected to be highest at the surface, and the stress decreases toward the center of the part or the center of the cross section. In this case, the ratio of the stress to the strength was higher below the surface than it was at the surface, so the crack initiated subsurface. It is important to note such an unusual situation, where the location of the ratio of the stress to the strength was highest at a subsurface position.
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Fig. 6
Fracture features of an induction-hardened shaft (1541 steel) after fatigue testing in rotary bending. A, B, fracture origins. Adapted from Ref 1, with annotations by W.T. Becker
Fig. 7
Stress and strength as a function of position in a cylindrical component loaded in torsion. Fracture initiation may be at either the surface or subsurface. Subsurface initiation depends strongly on the hardness profile from surface to center if loading is in bending or torsion.
Imagine a simpler case of a cylindrical component (Fig. 7). The surface of the part is shown along the left side of the graph, and the centerline is shown at right. For such a cylindrical component that is loaded in either bending or torsion, the stress will be highest at the outside surface, and at the centerline, or neutral axis, it will be nominally zero. A carburized or induction-hardened material is actually stronger at the surface layers where the stress is highest. The y-axis, instead of being the stress level, can be conceptually viewed as either the hardness or
some other measure of the strength of the component. If the part has a heavy case, then the strength follows the dashed line. In this case, high strength levels go in deep toward the core. At some point, the strength and hardness drop off to a lower level. In this situation, if the solid line represents the stress and the dashed line represents the strength, this part should not have a subsurface crack initiation. Everywhere, the stress is lower than the strength. If the case is too thin for the application in question, and the strength drops off as the dotted line shows, the stress is higher than the strength within a subsurface band, which allows subsurface crack initiation. This figure shows a powerful technique for specification of case depths, which has the potential to complement the usual experiential method of case depth specification. Anyone doing fracture analysis on a case-hardened part can also use this information to obtain an idea about the appropriateness of the hardening specification. It is important to realize that there is one other case where a subsurface initiation may occur that is not related to the heat treating or specification quality. Imagine the presence of a subsurface discontinuity, such as an inclusion, a void, or a tiny crack. Even if the steel has
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Stress and strength as a function of position in a cylindrical component loaded in torsion with subsurface discontinuities. Surface conditions may include: inadvertent decarburization, typically thin and may not be easy to find; deep case from induction or carburization; nitrided, thin case, often not more than 5–10 mils. Part of the case may be ground off in the finishing operations. The defect could be a faceted inclusion (nitride) in a low-ductility matrix. Nondeformable nitride causes stress concentration in the matrix.
or other rolling elements and races. Understanding that most loading geometries create the highest stresses at the surface also allows one to understand why decarburized layers can be so damaging. A decarburized layer is softer and lower in strength than the material with the desired nominal amount of carbon. Decarburization can even occur on carburized steel. Figure 9 shows a metallographic cross section of a piece of steel that is carburized and has quite a bit of retained austenite. Note the dark constituent at the surface (arrows). On carbon steels in the medium-carbon range, decarburization usually looks white, but here it looks dark due to the presence of pearlite. The decarburization affected the hardenability as well as the hardness in this case. There was not enough carbon to form martensite at the surface when the part was quenched in heat treatment. Very fine pearlite was formed instead. The pearlite structure is not as strong or fracture resistant as the martensite structure that is expected in the absence of the decarburization. Thus, this part could be more susceptible to fracture because of the decarburization during heat treatment. To complete the discussion of fracture, the previous is summarized by emphasizing that the macroscale features reveal the loading conditions. Fracture analysts must start with the macroscale, or the big picture. Many people start with the details, or the little picture, and move on to the big picture, but this can be a problem that facilitates mistakes on the part of the analyst. Microscale Fracture Features. Scanning electron microscopy can be used to reveal the microscale fracture features. Figure 10 shows a
Fig. 9
Fig. 10
constant strength all the way through its cross section, an inclusion that is big enough to locally increase the stress above the fracture strength of the part makes it possible to create a subsurface initiation (Fig. 8). Another potential cause of subsurface fracture initiation is contact loading. A well-lubricated bearing without any friction has the highest stresses in its subsurface layers. Thus, fracture initiation will be at a subsurface location where contact loads are the predominant source of stress. Inclusions can thus be very damaging in contact-stress applications, such as bearing balls
Fig. 8
Decarburization of carburized steel. Each small scale division is 2 mm.
Scanning electron micrograph of microvoid coalescence
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ductile crack path, or what is often called ductile dimples, microvoids, or microvoid coalescence (MVC). The MVC is a characteristic fracture morphology that indicates a component was subjected to stresses in excess of the nominal ultimate tensile strength of the component material. The MVC is often an indication that the heat treated steel did not have a gross problem with the heat treating that caused or contributed to an embrittlement problem. Brittle fracture features, including cleavage and (most) intergranular features, are often indicative of a heat treating problem. The arrows in Fig. 10 show the nonmetallic inclusions that initiated the void formation. Brittle fractures are often unexpected and occur suddenly without any prior warning. Ductile fracture by MVC is typically accompanied by prior plastic deformation, which gives advanced warning of the impending fracture event. This prior warning makes MVC the preferred mechanism if fracture occurs. While MVC is generally desirable, it can indicate that the material is too soft if a highstrength material is in question. The MVC can also reveal that the heat treater made a mistake, such as no heat treatment, despite the often preferred MVC fracture path. Classical microscale brittle crack paths are along the grain boundaries (intergranular), (Fig. 11) or cause the grains themselves to split (transgranular or cleavage). Many heat treating and other processing problems can cause undesirable intergranular cracking at the microscale. To have intergranular cracking, something may either cause a low-strength condition at the grain boundaries or cause the stresses at the grain boundaries to be higher than in the core of the grains. Refer to the previous concept regarding the relationship between the local stress and the local strength. One mechanism where intergranular fracture occurs is quench cracking. When a piece of steel is quenched to harden it, martensite will generally start forming near the surface, because that is where it cools off fastest. As anyone familiar with heat treating of steel knows, each deeper layer of grains will subsequently transform to martensite. On transforming to martensite, the material expands. As the material continues to cool, it contracts. So, there are grains that undergo expansion during transformation while the grains next to them contract. Could that create a shear stress at the grain-boundary
location? The situation is obviously more complicated than just described; obviously, the grainboundary strength is a function of temperature, which is rapidly changing, but it is helpful to think about what could be causing that intergranular crack and why quench cracks are generally intergranular. Note also that carburized steel often will have intergranular cracking in fatigue. Sometimes, even the best metallography cannot show any problem at the grain boundaries, that is, no grain-boundary carbides, oxides, nitrides, porosity, and so on. It is true that there are many heat treating problems that can facilitate, or be a factor in, intergranular cracking at lower stress levels than the part usually sustains. However, it is important to note that just because there is an intergranular crack, it does not mean there is surely a heat treating problem. Hydrogen embrittlement is often revealed in part by its intergranular crack habit, particularly in steels heat treated to high strength levels. Hydrogen embrittlement is not always purely intergranular; sometimes, there will be tiny, shallow microvoids on the grain-boundary surfaces. Hydrogen embrittlement could be the heat treater’s fault (if the hardness is too high), and it could be an issue with the plating (if records do not show proper baking); it could be an interaction between these two factors. The important point is not to confuse identification of the crack path with the cause. There is a difference between identification of the physical shape of the crack and the physical characteristics relating to the crack event. These do not automatically lead to the cause.
Fig. 11
Scanning electron micrograph of intergranular cracking
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The last microscale crack path is cleavage, and Fig. 12 shows a classical view. Ferrite cleaves readily at low temperatures. If a part is not supposed to have any ferrite in it and there is a large amount of cleavage, then that may be a clue to look carefully for ferrite during microstructural analysis. The classical way to recognize cleavage is the presence of patterns that look like riverbeds with multiple tributaries. The arrows in Fig. 12 show these river line features. In closing this section, fatigue microscale features in heat treated steels are often not very interesting or classical. The experienced analyst can recognize them, but they are difficult to describe. They rarely have the textbook striation features that are commonly shown for superalloys or aluminum alloys in published microfractographs. Striations may be visible in a low-carbon annealed steel, particularly in steels with ferrite as part of the microstructure. Beginners must be careful to distinguish pearlite platelets from striations. Pearlite stops at grain boundaries, and striations may cross grain boundaries. If there are questions, pearlite spacing can be examined on a cross section at a later time in an effort to distinguish one from the other. It is quite rare to find striations in any kind of hardened steel. Beach marks are often visible at the macroscale, but striations are very uncommon as microfractographic features. Summarizing the differences between macroand microscale features: The macrofeatures show the loading geometry. The microfeatures show the result of the microstructural interaction with the environment, and mechanical and chemical aspects may influence the way the crack interacts with the microstructure.
Fig. 12
Wear The original shape of the object shown in Fig. 13 was a gear with normal-shaped teeth. It is severely worn. No judgments can be made about the cause with this one image. Wear has many similarities to fracture and deformation. Wear is basically deformation and fracture going on at a microscale, and it can continue until the point that macroscale damage is present. Scanning electron microscopy (SEM) is very helpful in understanding how wear happens. Wear specialists have identified many different wear mechanisms. However, even without that specialized knowledge, the SEM can reveal useful information for diagnosis and prevention. In Fig. 14, there is smeared material. This solid steel has now flaked and smeared to the point that it is present as thin platelets, which are breaking off and allowing material loss. This is one example of a combination of deformation and fracture. Fretting is a common type of wear that is almost never related to heat treating or any problem with heat treating but rather is related to the geometry of the assembled parts. Fretting in steel will generally produce a reddish, iron oxide powder, and it roughens the surface. Figure 15 is an SEM image that just barely reveals the initiation of a crack. Fretting often produces a crack in an area that is thought to have low stress. Contact forces can cause surface damage due to the action of Hertzian stresses. Figure 16
Fig. 13 Scanning electron micrograph of cleavage cracking
Cross section of worn gear teeth. Approximate width of steel segment shown is 23 mm (0.9 in.).
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shows a bearing race and a bearing ball. The ball is being pressed into the race. This figure is a rough schematic, conceptually showing an exaggerated view of the most highly stressed area due to the elastic deformation of the two components. Without the load, there is essentially a point. Under load, the contact area becomes circular or elliptical; in other words, there is a contact footprint. The actual highstress location is just below the surface. In the presence of friction, the highest stresses are moved toward the surface. An important part of bearing wear failures is determining whether the crack initiated at the surface, showing the possibility of a lubrication issue, or if it was truly a subsurface initiation, in which case it may be a microstructure problem, an inclusion, or another anomaly. Bearings are generally loaded to very high stress levels. The quality of the steel and, in particular, minimization of inclusions are very important in this type of application. Another problem is that can cause surfaceinitiated cracks grinder burn, which can create high tensile stresses at the surface. Regardless of what the service load is, a stress field has been created with a very high tensile stress at the surface, the most undesirable location. To close this section on mechanical damage, it is often important to determine if there was a service problem, such as abuse, or misuse, such as using a screwdriver as a pry. Understanding loading geometries and related fragment shapes can shed light on this type of question. Figure 17 shows a steel bar with a threaded portion that is much smaller in diameter than the rest. Note that the crack is in the large-diameter
portion. Could this have been caused by the user? To answer the question, one would have to do a large amount of background information collection. The important point to realize is that the crack location is totally unexpected, and it is difficult to think of something that could have happened in service to create a weak spot at this hefty location. If it cannot be qualitatively demonstrated that somebody abused something, then it must be quantitatively demonstrated. This is often difficult. It is important to know what the material strength, fracture toughness, and other material properties were at the time of the damage. It is important to determine what the load or loading geometry was that caused the damage. One must
Fig. 15
Scanning electron micrograph of fatigue crack initiating on worn carbonitrided steel. Original magnification: approximately 4000 ·
Race
Bearing ball
Stress concentrations
Fig. 14
Scanning electron micrograph of a worn piece of hardened medium-carbon plate showing details of the wear mechanism
Fig. 16
Rough schematic of stresses in contact loading of a bearing ball on a race
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consider how the suspected loading geometry differs from the actual design intent. The analyst again must remember that just because it is broken does not mean it was abused. Even if it meets the specification, it does not prove that the user abused the product. Corrosion and Environmental Damage Figure 18 shows an example of a 300-series stainless steel that was probably not heat treated in the most ideal manner, since small precipitates can be seen along the grain boundaries. Note the crack location, which seems to be seeking the grain boundaries. To create a crack, there must be a load or a stress. There is corrosion present too, so this may be a stresscorrosion crack. Did the nonideal heat treating condition cause the crack? In this case, a chain of reasoning cannot directly link the specific
heat treating problem with the crack, because there is no information about the stress levels. Stress-corrosion cracking, by definition, requires a threshold level of corrosion and stress. In this case, there is no evidence that the crack was caused by bad heat treatment. This seemingly subtle distinction may be very important in the case of a catastrophic failure event. Figure 19 shows an example that involves another stainless steel weld. Welding is a kind of heat treatment, although not as controlled as an intentional heat treatment. An acid substance, polythionic acid, was in contact with the weldment. There are some cracks on one side of the weld, while the other side is free of cracks. What kind of cracks are these? The crack path is intergranular (Fig. 20). The micrograph in Fig. 21 was taken after an ASTM International test, and it shows ditching
Fig. 18
Grain-boundary precipitates in a 300-series stainless steel
Fig. 19
Stainless steel weldment
Fig. 17
Steel bar with crack in unexpected position. Originally shock loaded in compression. Threaded portion diameter is approximately 2 cm.
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characteristics. This indicates that whatever the heat sequence from the weld, it did create this condition that appears to have facilitated intergranular corrosion. However, steel would have been considered to pass the test for freedom from sensitization. Despite this, the thermal experience of the material can be directly linked to the form of the crack. Although this is not a heat treating example, this demonstrates the line of reasoning that is required to determine cause. It is interesting to note that one side cracked and the other side of the weld did not. It is likely that the side that did not crack was an L-series, a low-carbon series, specifically made to minimize the chance of cracking in weld heat-affected zones in stainless steels. The cracked material was probably not an L-series. This is an example of a classical
Fig. 20
Scanning electron micrograph of stainless steel weldment with intergranular cracking
Fig. 21
ASTM International sensitization test results showing ditching characteristics
damage mechanism, stress-corrosion cracking. When the regular 300-series stainless heated up, the chromium and carbon combined and precipitated along the grain boundaries, taking the chromium out of solution and making the material less corrosion resistant. Frequently, such situations lead to pit formation, which then allows the crack to propagate from the stress concentration at the pit. This is an example of a complex damage mechanism. Another commonly named damage mechanism is corrosion fatigue. It must be understood that a corroded part that broke due to repeated crack extension under load did not necessarily experience the mechanism called corrosion fatigue. Corrosion fatigue is a damage mechanism that is studied in the laboratory. The named mechanism is invoked when it can be demonstrated quantitatively that the crack is growing much faster under the same loading conditions than it would in the absence of the corrosive substance. In a real component out in the field, such as a heavy off-road vehicle application, it is very difficult to obtain an accurate service history day-by-day. Going back to published research data for standard test coupons and proving that a particular situation is or is not corrosion fatigue will likely prove very difficult. In Fig. 22, a shaft is
Fig. 22
Cracked shaft used in a corrosive environment. Diameter of the shaft is approximately 10 cm (4 in.).
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obviously corroded. The band highlighted by the arrows was covered with a layer of red rust. Figure 23 shows the crack created by the corrosion and stress combination. Note that there are also secondary cracks. The tip of this crack looks like a witch’s broom or the trails of a sparkler. This is not a normal fatigue crack. A normal fatigue crack does not branch out and have multiple tips. In this case, the corrosion definitely has some kind of significant impact on the damage mechanism. In many cases, examination of field failures is likely to leave a question mark regarding the quantitative evaluation of how much faster the crack is growing because of the corrosion. To demonstrate that corrosion is the cause or the fault of a failure, it is not enough simply to say the component was in a corrosive environment and it cracked because of that. It must be assured that the specific conditions, the specific material, and the specific process condition of the material were in the realm that has been demonstrated to be a problem for the damage mechanism invoked. For example, concentration of an aggressive substance, the threshold stress level, and the temperature may be required to be in a restricted range before a particular mechanism can be properly said to have been acting.
Fig. 23
Factors Contributing to Poor Response from Heat Treatment Raw Material Characteristics That Can Contribute to Poor Response from Heat Treatment What are the raw material characteristics that can contribute to poor heat treating outcome? One very important characteristic is composition. There will be a range of values for each type of atom that is specified for the grade in question, as well as for unspecified elements. A heat treater may receive material of an iron matrix that could, with the same name, have a very wide range of responses to the heat, heating rate, heating dwell time, cooling rate, and so on. Lean and rich alloy content can have a strong influence on whether or not quench cracks occur. Lean and rich compositions also strongly influence how readily the hardness specification is attained. For a hardness specification that is toward the upper limit of what can be reliably obtained for a particular grade, it may be difficult to meet the specification for a given lot if all the elements are on the lean end. If the part(s) crack on quenching, an important task in the troubleshooting process is to determine whether it is due to a lean or rich (more likely) alloy
Crack profile and adjacent secondary crack tip. Original magnification: 50 · . Inset is of a different secondary crack tip. Original magnification: 500 ·
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composition that is facilitating martensitic transformation. This must be differentiated from the situation where the major cause of the crack is that the heat treater quenched the part(s) more severely than usual. It is difficult to distinguish quench cracks due to extreme quench severity from those due to rich composition. However, it is easier to determine whether the crack has an expected geometry for a quench crack. In Fig. 24, the crack is at the section change, which is a prime location for a quench crack, as are sharp corners. When examining a part to determine if rich alloy/quench severity issues were at stake, examine the crack surfaces for traces of temper colors, including blues and browns. A dark matte gray or black surface may indicate the presence of an oxide-filled discontinuity that simply opened due to thermal stresses rather than a rich alloy/quench severity problem. If temper colors are visible at the portion of the crack surface nearest the part surface, it is likely that one or more of the following factors was present:
Quench was more severe than was appropriate Alloy was richer than usual
Section change was more severe (for example, the fillet radius was too sharp)
To confirm the suspicion of a “pure” quench severity/rich composition-related quench crack, it is advisable to confirm that the microscale crack path is intergranular (usually SEM is required). If a seam had been found at the quench-opened crack, it would not be correct to blame the heat treatment. Figure 25 shows an interesting crack. There is a very heavy oxide layer revealed by the cross section. The inset shows a higher-magnification view of the seam detail. There is a rounded particle that is totally covered by a heavy, rounded oxide layer. The crack surface does not have blue or brown temper colors but was found to be a dark charcoal black. This type of heavy oxide is unlikely to have happened between when the part was quenched (and presumably cracked) and when it exited the temper furnace. Another unusual feature of this crack is that it changed direction multiple times (large arrows show crack growth direction changes). The fact that these are not predominantly intergranular cracks deals the final blow to any theory stating that this crack was due to a heat treating problem. Some people wonder why it was not possible to see the seam before it went into the heat treating process. Many seams are tightly closed or smeared over until the part experiences the stresses of the heat treating operation. Limitations of nondestructive testing methodology may also play a role. In the case of a part with a quench crack in a location that would not be expected to have high stresses in quenching due to differential cooling rates, the experienced
Fig. 24
Quench crack with typical geometry
Fig. 25
Oxide layer along a seam most likely present in the raw material. Original magnification: 50 · . Inset original magnification: 200 ·
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analyst looks for some type of discontinuity. The discontinuity does not have to be very thick. This example was particularly heavy, but even a very thin oxide layer is enough to be the predominant cause of the crack. Macro- or microsegregation (otherwise known as banding) are other raw material characteristics that can interfere with the expected outcome of a heat treating process. Figure 26 is a longitudinal cross section from a mediumcarbon piece of steel. If the low-carbon layer is right at the surface, it may be difficult to meet a minimum hardness specification. Banding, or microsegregation, is not always bad. It can make it easier for the crack to grow in
Fig. 26
Longitudinal cross section showing microsegregation. Original steel segment shown is approximately
2 cm wide.
a particular direction, and that may be an advantage for a particular application. Samurai and Damascus swords from antiquity had basically banded microstructures with very desirable characteristics. While banding is not necessarily bad, it can cause some variation in the response to the heat treatment. If a material has intermittent coarse grains, it may be easier to form martensite in the large grains and pearlite, ferrite, or bainite in the surrounding fine grains. Can the heat treater create such a grain size distribution? This may be possible by overheating. However, it is also possible that the coarsening came from a subcritical amount of cold work stored in the material. Normalizing the steel prior to hardening may eliminate the nonuniform response to hardening, but the added cost of normalization is often objectionable (Fig. 27). Another type of raw material characteristic that can cause problems in some applications is heavy bands of stringers. Figure 28 shows a piece of steel that has long sulfide stringers in it, which can act like a seam. Figure 29 is of a wire product that was used in a coil spring. The material was subject to torsional loading. A longitudinal discontinuity in a material that is subject to torsional loading can create very high local shear stresses in the longitudinal or radial directions. This is one reason that seams are not allowed in critical applications for spring wires. Decarburization is not detected on the surface, but there is a heavy decarburized layer on either side of an inclusion of the shape characteristic of a seam. The heat treatment is unlikely to create such a varying thickness layer of oxide.
Fig. 27
A few coarse grains in the core of a fine-grained material that has been carburized are the only portion of the core able to form martensite. Original magnification: 100 ·
Fig. 28
Scanning electron micrograph of sulfide stringers in a piece of bar stock
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Component Characteristics Figure 30 shows a cast steel product that was quite uniformly carburized, except for the white script features. The pattern looks very similar to microsegregation between the dendrites and the matrix. It would be very difficult for the heat to overcome the initial segregation in the raw material. This component may not perform as well as a component with a uniform case. These
Fig. 29
Optical micrograph of an oxide-lined seam in a piece of steel wire
0
10
script features may make the case even more brittle than usual. The heat treater can control the carbon potential and the heating cycle. However, the heat treater cannot locally put fewer carbon atoms into the steel at the locations that already have too many. When performing failure analysis of steel components, analysts must
Fig. 30
20
Unusual crack shape Oxide filled crack Evidence of pre-existing discontinuities
Fig. 31
Thread root of a steel fastener. Original magnification: 100 ·
30
Cast steel after carburizing. Original magnification: 100 ·
40
Dark etching area possible indication of heat treating problem
Note multiple cracks in one thread root
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always be on the lookout for the whole history of the part that led up to the heat treating event. Forging discontinuities are another situation that could cause a problem. Figure 31 shows a cross section of a threaded fastener with a locally different compositional steel inclusion. Something unusual happened to create this feature. Other problems relating to poor heat treating or poor service outcome include many different types of design details. Other chapters in this book address these issues in greater detail. To summarize, some areas of common oversight include:
Materials selection Heat treating process selection Hardness level specification and range and position on part for hardness test Process details (batch or continuous oven, etc.) Heat treater’s familiarity with the size and complexity of the part and the quality level needed Distortion control
Fig. 32
Testing competence Hardness scale selection Frequency of part testing
Figure 32 shows two micrographs at the same magnification from the same component, 25 mm (1 in.) apart from each other. One of them is virtually all martensite. In the other location, there are wide grains of pearlite interspersed between the martensite grains. There was as much as 12 Rockwell C points difference between the two microstructures. The designer never indicated where to test the component! This is a typical design issue. Another common cause of disputes between purchasers and providers of heat treating services is that designers specify Rockwell tests when there is no way to perform anything but a Knoop or Vickers test. There are still many newly minted engineering prints with case depth specifications that are very unclear. There are standard methods for specifying carburized, induction-hardened, or carbonitrided case depths, and it is helpful to use an industry
Medium-carbon steel microstructures from the same component at two locations separated by approximately 25 mm (1 in.). Each small scale division is 5 mm.
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standard when possible. The automotive manufacturers have done a good job of providing a range of methods at a range of ease of testing. If a standard method is not being used, it is often difficult to determine the designer’s intent. The specifications for 400-series stainless steel can be particularly difficult for the average mechanical designer to write. Many designers specify 400-series stainless steel because they want stainless, but they want to be able to heat treat it to obtain higher hardness than standard 300-series annealed bar stock. For a number of the 400-series grades, one must determine in advance whether maximum strength or maximum corrosion resistance is desired. Two totally different heat treating processes attain those goals. The heat treater has no means of guessing which characteristic is required. What Are the Things That Can Go Wrong in the Heat Treating Process? This chapter has attempted to cover all the aspects for which the heat treater should not be blamed. What are the aspects for which the heat treater may or should bear responsibility? One approach is to say that heat treaters must take responsibility for those aspects that are specific to their equipment. These are details that
Fig. 33
only the heat treater could know. Design engineers specify materials and thus need to know how to specify the type of testing and evaluation required for application. A design engineer cannot know how fast a load of parts will be heated in a particular company’s individual furnace. The design engineer cannot be expected to know what type of fixturing may be necessary to maintain required distortion levels in the part. These are aspects the heat treaters must know. The heat treaters must know what load size can be treated in their own furnaces and how the load should be distributed. The heat treaters must understand the characteristics of the interactions between their equipment and the full range of part sizes and load sizes they are processing. Heat Treating Errors. Excessive heating rate, excessive time at temperature, and excessive temperature can lead to excessive distortion. Excessive temperature can cause problems with excessive autotempering. If a massive part is heated to 75, 100, or more degrees hotter than it needs to be to obtain the uniform austenite required before quench, then all the extra heat must be removed, which can make it difficult to create the desired martensite at all. Furthermore, the extra heat may act to partially temper the martensite that is present. Excessive
Undissolved ferrite and martensite in improperly specified and improperly induction-hardened medium-carbon steel part. Each small scale division is 5 mm.
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heating can cause trouble in obtaining minimum specified as-quenched hardness values. These low as-quenched hardness characteristics may or may not make a difference in the component performance. However, when as-quenched hardness tests are required, it is important to know that excessive temperature may be a cause. Inadequate Heating Rate, Inadequate Time at Temperature, and Inadequate Temperature. Figure 33 shows the microstructure of a part that was supposed to be induction hardened. There is some martensite. The light constituent is ferrite that never went into solution. The initial microstructure was probably a mixture of ferrite and pearlite. The pearlite transformed into martensite, but little of the ferrite did. On quenching, the material produced islands of martensite with islands of ferrite. The irregular shapes of the ferrite islands are classical undissolved ferrite. This is not a typical shape of ferrite grains formed on cooling. Because of the severity of the consequences, most of the common problems in heat treating are related to hardening. Annealing can also be done incorrectly. Figure 34 shows the microstructure of a steel that was supposed to be spheroidized annealed. Spheroidization is a process that may take 12, 15, or even 20 h at 600 to 700 C. Spheroidized annealed steel is generally quite expensive. Its applications are usually reserved for severe forming operations where
Fig. 34
Incompletely spheroidized annealed steel. Each small scale division is 2 mm.
the added ductility is a necessity. The annealing line management may be tempted to cut down the process time to save money. The minimum spheroidization that the material processor believes will work is what is often provided. In the case of varying incoming microstructures, the part may not be as ductile as usual. Examine the situation where the forming process creates a crack that was undetected, and then the part is heat treated for hardening. If the crack remains undetected, there is now a part with a discontinuity due to the spheroidizing being done poorly. This type of situation may be difficult to figure out, especially if the failure happens some years after the fact. Insufficient time or temperature could apply to formed parts requiring stress relief. Even at relatively low hardness values, excessive residual stresses can make the part sensitive to hydrogen embrittlement. A low-carbon steel part that has been heavily deformed and improperly stress relieved can crack after a very short service life or even while sitting on a storage shelf. Stress relief is often used on weldments, and if it is not done properly, fatigue cracks can initiate more readily at weld toes. Machined parts that are improperly stress relieved can distort or crack at a later time, because the stresses are higher than one may think. Cooled Too Fast. The part that is cooled too fast due to cold quenchant or excessive quenchant agitation may crack or suffer excessive distortion. Undesired microstructures may also result from excessively fast cooling. Bainite may be desired, but the process formed martensite. Cooled It Too Slowly or Cooled to the Wrong Temperature. This can be a result of a delay in moving the parts into the quench tank. Alternatively, the composition of the polymer quench tank may be improperly maintained. Slow cooling can be a problem, because the crack resistance of a microstructure with excess ferrite may be lower than a properly hardened and tempered martensitic structure. Improper Atmosphere. Decarburization can result from low carbon potential in the atmosphere surrounding the parts. Carburization can occur if there is too much carbon in the atmosphere. Retained austenite in undesirable amounts can also result from excessively rich carbon in the atmosphere. There are still heat treaters in business who believe that it is good to use some ammonia when carburizing, even if the
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customer did not ask for carbonitriding. These heat treaters may believe that the customer is happy to get the lower price due to a faster process that meets a surface hardness specification. The problem is that if the stress state requires a certain level of strength at a certain depth, the faster ammonia-enhanced process may be inadequate. Retained austenite in greater amounts than normally found in parts that are straight carburized can also be a problem. Porosity can also be created in the case with nitrogen atmospheres. This can be a problem if the surface hardening is desired for strength rather than just scratch resistance. Figure 35 shows an example of bad carburizing. There is a significant fraction of retained austenite in this case-hardened part, as well as a large, chunky, unusually shaped “puzzle piece” carbide. This feature could be a problem for some applications of carburized parts and may be the result of excessive carbon potential.
Concluding Comments For those who do failure analysis of heat treated steel parts, it is important to understand what microstructure is expected, given the heat treating process that is specified. For example, in
0
10
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a 1050 steel that is 100 mm in diameter and 100 mm long and water quenched, should it be martensite? There probably will be some martensite near the surface, but it may not be a very deep layer. The exact depth will depend on how hot it was heated prior to quench, the details of the quench tank design, and many other factors that may not be readily apparent. It is important to have a large amount of reasonably deep knowledge to be able to make a fair and correct determination of where there may have been a problem in the entire process of designing a part, procuring material, and making a component. This knowledge base includes failure analysis, fracture analysis, and microstructure analysis and interpretation in order to “read” the process history. Simply checking the hardness and the composition to see if they meet the specification is not failure analysis. Failure analysis includes a determination of the loading geometry and background information, at the very least, in addition to the basic certification conformance tests. Some people legitimately perform internal process failure analysis on a part that never left the door of the manufacturing plant. In failure analysis of a field return, even from a non-enduser assembly problem, it is important to do more than simply look at the composition and the hardness. The type of damage needs to be identified, as well as the possible sources of that
30
40
50
Fracture #2 has both large chunky and script carbides and large amounts of retained austenite.
Fig. 35
“Puzzle piece” carbide microstructure in carburized steel, possibly due to excessively high carbon potential. Each small scale division is 2 mm.
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type of damage. A review of the comprehensiveness of the design process may be in order. If recurrence prevention is a goal of the failure analysis, the damage specialist may lend some understanding to the design engineers to help them clarify the requirements of the component characteristics. Failure analysis can be very routine, or it can be extremely involved. This chapter has considered only a few categories of the analysis procedures and some of the reasoning involved in determining what went wrong and at what part of the life cycle the problem initiated. A decision must be made at the beginning of the analysis about how detailed the project will be. If there is a single component, especially, or a very limited number of failed parts, inadequate planning can leave inadequate specimen material for testing in the case of
unanswered questions at the end of the project. It is difficult to overemphasize the importance of spending enough time initially figuring out exactly what the goals of the failure analysis project are and how much detail is required.
ACKNOWLEDGMENTS The author thanks Mrs. W.T. Becker for permission to use the copyrighted material of William T. Becker in Fig. 6 to 8.
REFERENCE
1. D. Wulpi, Understanding How Components Fail, American Society for Metals, 1985
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Failure Analysis of Heat Treated Steel Components L.C.F. Canale, R.A. Mesquita, and G.E. Totten, editors, p 111-132 DOI: 10.1361/faht2008p111
Copyright © 2008 ASM International® All rights reserved. www.asminternational.org
General Aspects of Failure Analysis Waldek Wladimir Bose-Filho and Jose´ Ricardo Tarpani, Universidade de Sa˜o Paulo Marcelo Tadeu Milan, Instituto de Materiais Tecnolo´gicos do Brasil Ltda.
FAILURE ANALYSIS is the process of collecting, examining, and interpreting damage evidence. The objective is to understand the possible conditions leading to a failure and perhaps prevent similar failures in the future. A failure analysis should provide a welldocumented chain of evidence that either excludes or supports possible interpretation of the damage evidence. Clear-cut conclusions do not always occur, and the tendency of developing preconceived interpretations should be avoided. Various publications (e.g., Ref 1–6) describe the guidelines and methods of failure analysis, and this chapter briefly outlines some of the basic aspects of failure analysis. The first section describes some of the basic steps and major concerns in conducting a failure analysis. This is followed by a brief review of failure types from fracture, distortion, wear, and corrosion. Fracture is a common damage feature, because the vast majority of mechanical failures involve crack propagation—typically classified as ductile, brittle, and fatigue, as briefly described in more detail. Distortion, wear, and corrosion also can be important damage factors in failure analysis.
General Guidelines of Failure Analysis For a complete evaluation, the sequence of stages in the investigation and analysis of failure, as detailed in Ref 5, is as follows (Ref 2): 1. Collection of background data and selection of samples 2. Preliminary examination of the failed part 3. Nondestructive and mechanical testing 4. Selection, identification, preservation, and/ or cleaning of specimens
5. Macroscopic examination and analysis and photographic documentation 6. Microscopic examination and analysis 7. Selection, preparation, examination, and analysis of metallographic specimens 8. Determination of failure mechanism 9. Chemical analysis 10. Fracture mechanics analysis 11. Testing under simulated service conditions 12. Analysis of all the evidence, formulation of conclusions, and writing the report These stages or steps are briefly outlined as follows. Collection of Background Data and Selection of Samples. There are basically three fundamental principles to be carefully followed when collecting damage evidence from a fractured material (Ref 2):
Locate the origin(s) of the fracture. The whole fracture surface should be visually inspected to identify the location of the fracture-initiating site(s) and to isolate the areas in the region of crack initiation that will be most fruitful for further microanalysis. Where the size of the failed part permits, visual examination should be conducted with a low-magnification wide-field stereomicroscope having an oblique source of illumination (Ref 3). Do not put the mating pieces of a fracture back together, except with considerable care and protection. Protection of the surfaces is particularly important if electron microscopic examination is to be part of the procedure (Ref 2). Appropriate packaging of failed components for shipping is equally important. Wrapping them directly into a plastic bag, or placing pieces directly into a plastic bottle or container, can introduce unwanted hydrocarbon contaminants.
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Fingerprints on the failed surfaces can also introduce contamination (Ref 4); Do not conduct a destructive testing without considerable thought. Alterations such as cutting, drilling, and grinding can ruin an investigation if performed prematurely. Destructive testing must be performed only after all possible information has been extracted from the part in the original condition and after all significant features have been carefully documented by photography (Ref 2).
Preliminary Examination of the Failed Part. In addition to locating the failure origin, visual analysis is necessary to reveal stress concentrations, material imperfections, presence of surface coatings, case-hardened regions, welds, and other structural details that contribute to cracking. A careful macroexamination is necessary to characterize the condition of the fracture surface so that the subsequent microexamination strategy can be determined. Corrodents often do not penetrate the crack tip, and this region remains relatively clean. The visual macroanalysis will often reveal secondary cracks that have propagated only partially through a cracked member. These part-through cracks can be opened in the laboratory and are often in much better condition than the main fracture (Ref 3). Nondestructive and Mechanical Testing. A wide variety of nondestructive testing is available, including dye penetrant, ultrasonics, x-ray, and eddy current, which can help in the failure analysis task in order to unveil even subtle and/or internal defects in a part. Mechanical property tests are also ready to use, ranging from a sample hardness test to elevatedtemperature tensile and impact testing. These tests are often used to determine if degradation is related to fabrication or to the service environment. Sometimes, a standard test can be adapted to simulate manufacturing or in-service conditions more closely (Ref 4). Selection, Identification, Preservation, and/or Cleaning of Specimens. Unless a fracture is evaluated immediately after it is produced, it should be preserved as soon as possible to prevent attack from the environment. The best way to preserve a fracture is to dry it with a gentle stream of dry compressed air, then store it in a desiccator, a vacuum storage vessel, or a sealed plastic bag containing a desiccant. However, such isolation of the fracture is often not practical. Therefore, corrosion-preventive
surface coatings must be used to inhibit oxidation and corrosion of the fracture surface. The primary disadvantage of using these surface coatings is that fracture surface debris, which often provides clues to the cause of fracture, may be displaced during removal of the coating. However, it is still possible to recover the surface debris from the solvent used to remove these surface coatings by filtering the spent solvent and capturing the residue. In regard to cleaning techniques, fracture surfaces exposed to various environments generally contain unwanted surface debris, corrosion or oxidation products, and accumulated artifacts that must be removed before meaningful fractography can be performed. Before any cleaning procedures begin, the fracture surface should be surveyed with a low-power stereobinocular microscope, and the results should be documented with appropriate sketches or photographs. Low-power microscope viewing will also establish the severity of the cleaning problem and should also be used to monitor the effectiveness of each subsequent cleaning step. It is important to emphasize that the debris and deposits on the fracture surface can contain information that is vital to understanding the cause of fracture. The most common techniques for cleaning fracture surfaces, in order of increasing aggressiveness, are (Ref 3):
Dry air blast or soft organic-fiber brush cleaning Replica stripping Organic-solvent cleaning Water-based detergent cleaning Cathodic cleaning Chemical-etch cleaning
Macroscopic Examination and Analysis and Photographic Documentation. More often than not, the investigation starts with a low-magnification, if any, observation of the failed part. This visual examination can often quickly answer questions such as: What was the mode of failure? Did it crack, or was there a uniform or pitting corrosion failure? Did the protective oxide film break down? Were the welds visibly contaminated? A variable magnification stereoscope equipped with a ring light and directional fiberoptic lighting is a powerful tool for macroscopic visual examination. Contemporary stereoscopes can operate over a range of 2.5 to 50 · (Ref 4). Microscopic Examination and Analysis. Once the area of interest is isolated, a smaller
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portion can be cut from the sample and mounted for metallographic polishing and microscopic examination. The microstructure of specimens may be enhanced by a wide variety of metallographic techniques that include, for example, heat tinting, stain etching, anodizing, and illumination by bright-field and polarizing light. Optical microscopic examination generally begins at 50 · magnification and continues through 1000 · or even 1500 · . Higher levels are best supplemented by differential interference contrast lighting, which allows theoretical resolution of features as fine as one-third of a micrometer. Features that are important to recognize include the uniformity and size of the grain structure, the size distribution and shape of intermetallic particles, and inclusions. Scanning electron microscopy (SEM) is most useful where extreme depth of focus and high magnifications are needed. Fractures generally are complex, undulating surfaces that are difficult to image, and an optical microscope can only focus on a very narrow region because of the very shallow depth of field. However, the SEM excels at imaging fracture surfaces, and it can be operated in many different modes. The most common mode is secondary electron imaging, which provides a detailed, high-depth-focus image that is easy to interpret. Backscattered “Z” contrast is used to identify regions of impurities within a matrix. High-atomic-number species produce a light appearance, whereas low-atomic-number species create a darker appearance. The topographic backscattered mode enhances the surface topography of the sample and accentuates height or elevation differences on a fracture surface. The characteristic x-rays can be detected and analyzed according to their energy. This is called energy-dispersive x-ray analysis. The x-ray wavelength corresponds to the presence of a specific element, and its amplitude corresponds to the quantity of such element. This technique allows quantitative characterization of elements within a given phase. Bulk chemistry is typically analyzed during failure analysis to verify conformance with industry-accepted chemical limits. In the case of reactive metals, light elements can embrittle them due to improper processing or service conditions (Ref 4). Selection, Preparation, Examination, and Analysis of Metallographic Specimens. One of the worst things that can happen to the sample is inadequate handling, examination, or packaging. It is imperative that the sample remains in
an undisturbed state prior to analysis, because the culprit is often found in minute surface features or traces of impurities. Fracture surfaces must remain untouched so that high-magnification images can accurately determine the failure mode. The sample must be removed carefully. Important evidence can be destroyed by overheating or by allowing adjacent fracture surfaces to fret or rub together during sectioning. The ideal method would be to unbolt the component or to provide adequate support so that a slowspeed saw can be used to cut out the component. However, sawing lubricants can mask or destroy residual chemicals or elements on the failed surface, so precautions become extremely necessary. If the component has failed in the middle of a large area, more aggressive cutting/ sectioning techniques may be warranted, but keep a good distance from the failed region (Ref 4). Determination of Failure Mechanism (with Adapted Text from Ref 7). A thorough investigation should ensure that all damage is found and documented, because multiple modes and mechanisms may be present in most realworld failure analyses. It is also important to recognize that many unique mechanisms may be driven by more than one environmental factor, such as stress, temperature, corrosion, wear, radiation, or electrical factors. The term failure mechanism, or damage mechanism, is meant to convey the specific series of events that describe both how the damage was incurred and the resulting consequences. Examples of damage mechanisms include high-temperature creep, hydrogen embrittlement, stress-corrosion cracking, and sulfidation. A failure or damage mechanism describes how damage came to be present. This definition of failure mechanism also should not be confused with the description of the physical characteristics of damage observed. For example, intergranular fracture, buckling, transgranular beach marks, and pits can all be thought of as damage modes. The term damage mode or failure mode is best used to describe what damage is present. Much confusion has occurred because of the tendency of engineers to use the terms mechanism and mode interchangeably; in doing so, it is unclear that two distinct characteristics need to be assessed. Sometimes this occurs because, within a given system, the same wording is used to describe both the failure mode and mechanism. For example, pitting
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describes a damage mode because the surface of a material is pitted. In certain systems, pitting is also a possible damage mechanism. In boiler tubing, for example, a pitting damage mechanism describes a specific localized corrosion mechanism where pits form through dissolution of metal either from low-pH or high-oxygen conditions. The metal under the pit surfaces is unaffected. In this system, pitting is a specific damage mechanism, but many other damage mechanisms also result in a pitting damage mode in boiler tubing, including hydrogen damage, phosphate corrosion, and caustic gouging. It is helpful to be as specific as possible in differentiating damage mechanisms in a system. For example, fatigue is often identified as both a damage mode and a damage mechanism. A fatigue damage mode is the observable damage that occurs under fatigue loading cycles (e.g., the presence of beach marks). Classifying fatigue as a damage mechanism is not necessarily complete because it does not point to the specific environment that results in a fatigue damage mode. Instead, specific mechanisms that can result in a fatigue damage mode must be examined. Examples include corrosion fatigue, thermomechanical fatigue, creep-fatigue interaction, and mechanical fatigue. Determination of damage mechanisms starts by characterizing the component(s) being examined. It is impossible to know what is different about a failure without first understanding what is expected from unfailed components. In general, the analyst should obtain as much information as possible about a part and its background during the course of an investigation. Some key questions worth evaluating include:
What was the part supposed to do? How was it supposed to work? How was the part made? What processes were involved in its manufacture (e.g., forming, joining, and heat treatment)? What properties were expected at the time of manufacture? What were the specified dimensions and tolerances for the as-manufactured part? How was the part installed? To what service environment(s) was the part exposed? Typical environments to examine include operating temperatures, stresses (steady state or slowly rising and cyclic), oxidizing/corrosive environments, and wear
environments. What properties were required during service? How were properties expected to change from service exposure? How was the part inspected during service intervals? What information was found during these inspections? What material characteristics were specified for the part (e.g., composition, strength, hardness, impact, and stress-rupture properties)? What specifications, industry standards, and contracts govern these properties? What were the various ways the part could fail?
The last item is a key question to repeatedly ask throughout a failure investigation. The list of various damage mechanisms by which a part can fail can be narrowed down through two basic concepts (Ref 7). Limiting conditions that refine the scope of explanations for observed damage can be defined by using the following two rules of thumb:
When the impossible is eliminated, whatever remains, however improbable, must be considered (Sherlock Holmes rule). When two or more explanations exist for a sequence of events, the simple explanation is more likely to be the correct one (Occam’s razor).
Chemical Analysis. In a failure investigation, routine analysis of the material is usually recommended. There are two main categories of chemical analysis that are often used by failure analysts:
Bulk composition evaluation: often performed in order to determine whether the correct alloy was used in the subject component Microchemical analysis: to find evidence of contamination, to evaluate the composition of microphases revealed on a metallographic specimen, or to evaluate corrosion products
Often, chemical analysis is done last, because an analysis usually involves destroying a certain amount of material. There are instances where the wrong material was used, under which conditions the material may be the major cause of failure. In many cases, however, the difficulties are caused by factors other than material composition. Extreme care must be used in interpretation of chemical analysis work performed as part of a failure investigation. Minor deviations from
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specified composition must not be interpreted as the sole cause of a failure, without much additional supporting evidence. In most instances, slight deviations from specified compositions are not likely to be of major importance in failure analysis. However, small deviations in aluminum content can lead to strain aging in steel, and small quantities of impurities can lead to temper embrittlement. In specific investigations, particularly where corrosion and stress corrosion are involved, chemical analysis of any deposit, scale, or corrosion product, or a substance with which the affected material has been in contact, is required to assist in establishing the primary cause of failure. Where analysis shows that the content of a particular element is slightly greater than that required in the specifications, it should not be inferred that such deviation is responsible for the failure. Often, it is doubtful whether such a deviation has played even a contributory part in the failure. For example, sulfur and phosphorus in structural steels are limited to 0.04% in many specifications, but rarely can a failure in service be attributed to sulfur content slightly in excess of 0.04%. Within limits, the distribution of the microstructural constituents in a material is of more importance than their exact proportions. An analysis (except a spectrographic analysis restricted to a limited region of the surface) is usually made on drillings representing a considerable volume of material and therefore provides no indication of possible local deviation due to segregation and similar effects. Also, certain gaseous elements, or interstitials, normally not reported in a chemical analysis, have profound effects on the mechanical properties of metals. In steel, for example, the effects of oxygen, nitrogen, and hydrogen are of major importance. Oxygen and nitrogen may give rise to strain aging and quench aging. Hydrogen may induce brittleness, particularly when absorbed during welding, cathodic cleaning, electroplating, or pickling. Hydrogen is also responsible for the characteristic halos or fisheyes on the fracture surfaces of welds in steels, in which instance the presence of hydrogen often is due to the use of damp electrodes. These halos are indications of local rupture that has taken place under the bursting microstresses induced by the molecular hydrogen, which diffuses through the metal in the atomic state and collects under pressure in pores and other discontinuities. Various effects due to gas
absorption are found in other metals and alloys. For example, excessive levels of nitrogen in superalloys can lead to brittle nitride phases that cause failures of highly stressed parts. Various analytical techniques can be used to determine elemental concentrations and to identify compounds in alloys, bulky deposits, and samples of environmental fluids, lubricants, and suspensions. Semiquantitative emission spectrography, spectrophotometry, and atomicabsorption spectroscopy can be used to determine dissolved metals (as in analysis of an alloy), with wet chemical methods used where greater accuracy is needed to determine the concentration of metals. Combustion methods ordinarily are used for determining the concentration of carbon, sulfur, nitrogen, hydrogen, and oxygen. Wet chemical analysis methods may be employed for determining the presence and concentration of anions such as Cl , NO3 , and S . These methods are very sensitive. X-ray diffraction identifies crystalline compounds either on the metal surface or as a mass of particles and can be used to analyze corrosion products and other surface deposits. Minor and trace elements capable of being dissolved can be determined by atomic-absorption spectroscopy of the solution. X-ray fluorescence spectrography can be used to analyze both crystalline and amorphous solids, as well as liquids and gases. Stress Analysis and Fracture Mechanics Analysis. When confronted with a cracked, fractured, or deformed component, the failure analyst will usually seek to answer some basic questions:
Were the loads and stresses encountered by the part at the level anticipated during design? Or did some unexpected condition(s) contribute to the failure? Was the material in the area of the cracking or deformation capable of meeting the conditions anticipated during design? Was there some deficiency or discontinuity that contributed to the failure, or was there a local stress raiser at the critical location? Was this taken into account by the designer?
In general, there are two types of conditions that may lead to structural failure:
Net-section instability, where the overall structural cross section can no longer support the applied load
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